Patent Application
20060013719
1. Field of the Invention
The present invention relates to a wear-resistant sintered aluminum alloy with high strength that is suitable for various kinds of sliding parts such as connecting rods, pistons and the like, and a method of manufacturing thereof. More particularly, the invention concerns a wear-resistant sintered aluminum alloy with high strength that is improved in tensile strength and elongation as well as wear resistance, and a manufacturing method thereof.
2. Related Art
Regarding the aluminum sintered parts manufactured with the use of a powder-metallurgical method, there has been an increasing demand in recent years, since they are not only light in weight but also possible to possess preferable properties that cannot be obtained with cast materials, such as strength, wear resistance and the like. Specifically, in a case of wrought alloy containing a large amount of silicon, only alloy having metallographic structure in which primary silicon grains are coarse is obtained. In contrast, in a case of sintered aluminum alloy, it has been accomplished to obtain a sintered aluminum alloy which has metallographic structure that Al—Si alloy phase having fine primary silicon grains dispersed therein and aluminum solid solution phase having no primary silicon grains are dispersed in spots, and which is excellent in strength and wear resistance (ref. publications of Japanese Laid-Open Patent applications, JPA-H04-365382, JPA-H07-197168, JPA-H07-197163, and JPA-H07-224341). These sintered aluminum alloys are excellent in wear resistant. However, they are to an extent of 360 MPa or so in terms of the strength even when they are subjected to forging and heat treatment, and the application of them is limited and a sintered aluminum alloy with a higher level of strength has been therefore expected to be produced.
In short, it is not such a material that exhibits a high level of property in terms of both of the tensile strength and the elongation.
With the above problems in mind, it is therefore an object of the present invention to provide a novel sintered aluminum alloy having a wear resistance and simultaneously having a higher tensile strength and a high elongation, and a method of manufacturing the same.
It is also an object of the present invention to provide a manufacturing method of a sintered aluminum alloy, wherein the zinc content of the sintered aluminum alloy after the sintering is not fluctuate to achieve a constant mechanical strength and realize stable mass production thereof.
In order to achieve the above-mentioned object, a wear-resistant sintered aluminum alloy having high strength, according to one aspect of the present invention, comprises, by mass: 3.0 to 10% zinc; 0.5 to 5.0% magnesium; 0.5 to 5.0% copper; 0.1 to 10% hard particles; an inevitable amount of impurities; and aluminum, and having a metallographic structure comprising: an aluminum alloy matrix in which the hard particles dispersed; and an intermetallic compound phase being dispersedly precipitated in the aluminum alloy matrix.
A method of manufacturing a wear-resistant sintered aluminum alloy with high strength, according to one aspect of the invention comprises: preparing a raw material powder comprising, by mass: 3.0 to 10% zinc; 0.5 to 5.0% magnesium; 0.5 to 5.0% copper; 0.1 to 10% hard particles; inevitable amount of impurities; and the balance aluminum, by using: an aluminum powder having a particle size of 140 microns or less; a powder for the hard particles, having a particle size of 113 microns or less; and one of combination of simple metal powders, combination of binary alloy powders and combination of a simple metal powder and a binary alloy powder, containing zinc, magnesium and copper and having a particle size of 74 microns or less; pressing the raw material powder in a die at a compacting pressure of 200 MPa or more to form a compact having a predetermined shape; sintering the compact in a non-oxidizing atmosphere in such a manner as to heat the compact from 400 degrees C. to a sintering temperature of 590 to 610 degrees C. at an temperature-elevating rate of 10 degrees C./min or more and keep the sintering temperature for 10 minutes or more, before cooling the sintered compact to a room temperature; and subjecting the compact after the sintering, to heat treatment comprising: heating the compact at a temperature of 460 to 490 degrees C. and water-quenching so as to dissolve a precipitation phase in the aluminum base of the compact to produce solid solution; and keeping the temperature in a range of 110 to 200 degrees C. for 2 to 28 hours to produce a precipitation phase from the solid solution.
A method of manufacturing a wear-resistant sintered aluminum alloy with high strength, according to another aspect of the invention, comprises: preparing a raw material powder comprising, by mass: 3.0 to 10% zinc; 0.5 to 5.0% magnesium; 0.5 to 5.0% copper; 0.1 to 10% hard particles; inevitable amount of impurities; and the balance aluminum, by using: a simple aluminum powder of at least 15 mass % of the raw material powder; an aluminum alloy powder containing the whole of zinc which the raw material powder comprises; and a powder for the hard particles at 0.1 to 10 mass % of the raw material powder; pressing the raw material powder in a die at a compacting pressure of 200 MPa or more to form a compact having a predetermined shape; sintering the compact in a non-oxidizing atmosphere at a sintering temperature of 580 to 610 degrees C. for 10 minutes or more, before cooling the sintered compact to a room temperature; and subjecting the compact after the sintering, to heat treatment comprising: heating the compact at a temperature of 460 to 490 degrees C. and water-quenching so as to dissolve a precipitation phase in the aluminum base of the compact to produce solid solution; and keeping the temperature in a range of 110 to 200 degrees C. for 2 to 28 hours to produce a precipitation phase from the solid solution.
In accordance with the above construction, the sintered aluminum alloy of the present invention is excellent to have high tensile strength and elongation as well as high wear resistance. In the manufacturing method of a sintered aluminum alloy of the present invention, the tensile strength and the elongation are especially improved for wear resistant sintered aluminum alloys, to enable application to various kinds of sliding members used in vehicles and realize various sliding members of small weight.
The features and advantages of the manufacturing method according to the present invention over the conventional art will be more clearly understood from the following description of the preferred embodiments of the present invention.
The inventors of the present application has developed and proposed in Japanese Patent Applications of No. 2003-345001 and No. 2004-206957, a manufacturing method of an aluminum alloy having composition of ASM (American Society for Metals) 7xxx series that are known as extra super duralumin, by a powder metallurgical process. In this application, a wear-resistant sintered aluminum alloy simultaneously having a high strength is accomplished, based on the above sintered alloy, by the addition of hard particles in the composition. Moreover, the manufacturing method is further improved to prevent vaporization of zinc and variation of the zinc content by incorporating the whole amount of zinc in the form of aluminum alloy, and to prevent decrease of compressibility of the raw material powder mixture, by use of at least 15 mass % of a simple aluminum powder in combination with the aluminum alloy powder as raw material powders.
Hereinafter, the embodiments of the present invention will be described, explaining every component and step of the manufacturing method in detail. It is to be noted that, in the description of the present application, Al, Zn, Mg and the like are symbols of the elements used, and that the term, “aluminum part(s)”, should be read as meaning “aluminum-based part(s)” or part(s) being composed mainly of aluminum and possibly containing small amounts of other elements. Moreover, “sintering temperature” means the maximum temperature at which the compact is sintered, and “sintering time” means the time period during which the temperature is in the range of the sintering temperature.
(1) Raw Material Powder Blending Step
In this step, a raw material powdery mixture to be compacted is prepared by blending the respective powdered raw materials as to which the details are described below.
(1)-1 Ingredient Composition
<Zinc>
Zinc, together with magnesium, is precipitated in the aluminum matrix in the form of MgZn2 (η-phase) or Al2Mg3Zn3 (T-phase) to work to make an increase in the strength. Also, zinc, when the temperature is elevated for sintering, is molten to become a liquid phase and it wets the surface of the aluminum particles to eliminate the oxide layer thereon, and it is diffused into the aluminum matrix to also act to accelerate the bonding of the aluminum particles resulting from diffusion of them to each other due to such diffusion of zinc. If the content of Zn is below 3 mass %, it is difficult to sufficiently exhibit the above-described works, with the result that the effect of making the enhancement in the strength becomes poor. On the other hand, if the content is beyond 10 mass %, the amount of Zn in the sintered mass or the amount of Zn-based eutectoid liquid phase becomes excessively large, with the result that it becomes impossible to maintain the shape of compact as is. In addition, the portion where Zn is excessive or the diffusion of Zn into the Al base is insufficient remains in the form of a Zn-rich phase. Further, Zn volatilizes from inside the alloy and in consequence contaminates the interior of the furnace and is deposited there. Accordingly, the content of Zn is preferred to range from 3 to 10 mass %.
<Magnesium>
Magnesium forms the above-described precipitation compound together with zinc to contribute to enhancing the strength. Also, Mg is also low in melting point, and in the course where the temperature is elevated for performing the sintering, it produces a liquid phase to eliminate the oxide layer to work to accelerate the progress of the sintering. If the content of Mg is below 0.5 mass %, that makes the above-described effect poor, and, if it is over 5.0 mass %, that increases the amount of liquid phase to become excessively large, resulting in that it becomes impossible to maintain the shape of compact as is. Accordingly, the content of Mg is preferred to range from 0.5 to 5.0 mass %.
<Copper>
Copper is dissolved in the aluminum matrix to form solid solution and precipitate a compound of CuAl2 (θ-phase), thereby contributing to enhancing the strength. It also generates a liquid phase, when performing the sintering step, and works to accelerate the progress of the sintering. Regarding the content of Cu, if it is below 0.5 mass %, that work is not sufficiently attained, and, if it exceeds 5.0 mass %, copper forms an unnecessary Cu—Zn alloy phase with zinc, which is precipitated large along the grain boundary to cause the decrease in the strength and elongation. Therefore, the amount of Cu is preferred to range from 0.5 to 5.0 mass %.
<Hard Particles>
The metallographic structure of the aluminum alloy matrix composed of the above-described components, having no hard particles, exhibits excellent mechanical properties which are equal to those of the general steel materials so that the tensile strength is 500 MPa or more and the elongation is 4% or more, when making proper the conditions for compacting, sintering, forging and heat treatment.
In general, incorporation of a hard phase into an alloy matrix leads to decrease in strength and elongation of the alloy. However, since the aluminum alloy matrix as the base is made into alloy with the above-described elements to impart strength, extremely high strength and elongation in comparison with those of the conventional wear-resistant sintered aluminum-silicon alloy are possibly exhibited beyond a slight decrease in strength and elongation by the addition of hard particles. Moreover, in the present invention, it is advantageous in that the kind and amount of the hard particle dispersed can be easily changed in accordance with the sliding conditions (particularly, the sliding counterpart). Specifically, the hard particles dispersed in the conventional wear-resistant sintered aluminum-silicon alloy are primary silicon crystal, and it tends to increase the coefficient of friction when the sliding counterpart is made of ferrous material, due to the affinity between iron and silicon. In contrast, in the sintered aluminum alloy of the present invention, it is possible to reduce the coefficient of friction and improve wear resistance, by selecting a kind of hard particles having low affinity to iron, such as chromium boride or the like. The hard particles that can be used in the present invention includes silicon carbide, chromium boride, boron carbide and the like.
When the content of the hard particles in the sintered aluminum alloy is 0.1 mass % or more, the effect of improving wear resistance is distinguished, and, if it is over 10 mass %, the strength and elongation are remarkably reduced. Therefore, the content of hard particles is preferably in a range of 0.1 to 10 mass %. If the hardness of hard particles is insufficient, the hard particles themselves cause plastic flow, resulting in decrease of wear resistance. Accordingly, the Vickers hardness of hard particles is preferably 600 Hv or more, and more preferably 1,000 Hv or more, especially when an Al—Zn alloy powder is used as a raw material.
<Sn, Bi, In >
Tin, bismuth and indium are low in melting point and generate a liquid phase in the sintered mass, respectively. As a result, they wet the surface of the aluminum particles and eliminate the oxide layer from the surface of the aluminum particles, to accelerate the progress of the sintering between the Al powder particles without solution in aluminum. In addition, due to the surface tension of liquid phase, the liquid phases cause shrinkage, which works to contribute to densifying the resulting mass. Therefore, it would be preferable if using the above elements as a sintering auxiliary agent together with the above-described Zn, Mg and Cu. When the length of the term during which the liquid phase exists increases, the densifying attributable to the liquid phase progresses further. Therefore, if the liquid phase generates at an early stage of the sintering step so that the liquid phase continues to exist during the almost entire step of sintering, the densifying effect becomes great. In this view point, Sn (the melting point: 232° C.), Bi (the melting point: 271° C.), and In (the melting point: 155.4° C.) are very suitable, because they have a low melting point and they are hardly dissolved into the main component, Al. Moreover, the liquid phase of these low-melting-point metal components covers the surface of the simple zinc powder or the zinc alloy powder to prevent vaporization of zinc and fluctuation of zinc content in the resulting sintered alloy.
Sn, Bi and In which are auxiliarily used as the sintering aid agent may be used in the form of simple metal powder. If using these elements as the main components and forming an eutectic compound which would cause the production of an eutectic liquid phase comprising those main components, its melting point is much lower than that in the case of single substance. Therefore, making into that eutectic compound is further preferable. This eutectic liquid phase may be the one that is made by combining the main component (Sn, Bi, In) and another element, or the one which is made by combining the main component and an intermetallic compound that comprises the main component and another element. Moreover, there is also a compound having a line of eutectic reaction which can be found in a part of the monotactic compounds, and it is also possible to use such a monotactic compound causing the production of a eutectic liquid phase that comprises Sn, Bi or In. As the elements which form the eutectic liquid phase like that with Sn, there are Ag, Au, Ce, Cu, La, Li, Mg, Pb, Pt, Tl, Zn and the like. As the elements which form the eutectic liquid phase like that with Bi there are Ag, Au, Ca, Cd, Ce, Co, Cu, Ga, K, Li, Mg, Mn, Na, Pb, Rh, S, Se, Sn, Te, Tl, Zn and the like. As the elements which form such a eutectic liquid phase as described above with In, there are Ag, Au, Ca, Cd, Cu, Ga, Sb, Te, Zn and the like. Although these respective groups of elements are an example of simple two-elemental or binary system, the same effect can be obtained even in a case of a three-elemental or ternary system, a four-elemental or quaternary system or more-elemental system, as long as the resulting eutectic liquid phase similarly has Sn, Bi or In as the main component and has a composition causing the production of a eutectic liquid phase that comprises the main component. However, regarding Pb and Cd of the above elements, although these elements also cause the production of a eutectic liquid phase with Sn, Bi or In, it is preferable to abstain the use of them from the standpoint of toxicity.
With the above-described standpoint also being taken into consideration, as a multi-elemental system of eutectic alloy that comprises Sn, Bi or In as the main component, a lead-free solder the development of which has in recent years been urged can be preferably used. As the lead-free solder, ones of Sn—Zn system, Sn—Bi system, Sn—Zn—Bi system, Sn—Ag—Bi system or the like can be given, and lead-free solders prepared by adding to the above system a small amount of metal element such as In, Cu, Ni, Sb, Ga, Ge or the like has been proposed. A part of them has actually been put into practical use, and it is preferable to use such lead-free solders that are commercially available, since this it is easy to obtain.
The sintering auxiliary agent, when added 0.01 mass % or more, exhibits a remarkable densifying effect. However, if used in large amount, Sn, Bi and In become precipitated at the grain boundary to cause the decrease in the strength, since they are not dissolved with Al. Therefore, the use of them should be limited to 0.5 mass % or less at the most. Adding in an amount of 0.5 mass % or more results in that the decrease in the strength due to the precipitation of the Sn, Bi and In at the grain boundary becomes larger in degree than the above-described effect of densification due to the shrinkage of the liquid phase, resulting in more decrease in strength.
(1)-2 Form of Powder
A. In a Case of Using a Simple Zinc Powder
Regarding the above-described Zn, Mg and Cu, no inconvenience occurs when they are added to the zinc powder, in any case of using simple element powder, alloy powder of two or more kinds of these elements, or a powdery mixture of them. However, in order to cause the above-described works uniformly in the base, it is necessary to disperse the respective ingredient elements uniformly in the matrix. For this reason, it is recommended that those ingredient elements, as later described, are added in the form of fine powder whose particle size is 200 meshes (74 microns) or less. In a case where they are added like that, the simple element powder or alloy powder is melted when the temperature is elevated during sintering and becomes a liquid phase to wet the surface of the aluminum powder to eliminate the oxide layer thereon. They are then diffused into the aluminum matrix and, in addition, accelerate the bond between the aluminum powder particles due to such diffusion. If the particle size of the simple element powder or alloy powders exceeds 200 meshes, local segregation occurs to inhibit uniform diffusion of components.
In contrast, if the simple aluminum powder is also so fine as the above, the flowability of the raw material powdery mixture is reduced. Therefore, as for the aluminum powder, it is preferable to use in a larger particle size than the powders for above described elements. However, if the particle size exceeds 100 meshes (140 microns), the above-described components are difficult to diffuse into the aluminum particles, resulting in segregation.
B. In a Case of Using Aluminum Alloy Powder Containing the Whole Amount of Zinc
Zinc is an element that is likely to volatilize at a high temperature. Therefore, if Zn is added in the form of aluminum alloy powder by alloying the whole amount of Zn with aluminum, the amount of Zn that remaining through the volatilization of Zn becomes more stable than that in a case where Zn is added as simple zinc powder. As a result of this, the degree of fluctuation among the products becomes small.
However, incorporation of Zn causes a hardening in the aluminum alloy powder to decrease the compressibility of the powder. Accordingly, if Zn is made into alloy with the whole amount of aluminum, the compressibility of the raw material powder is decreased. Therefore, it is necessary to limit the use of aluminum alloy powder containing zinc to only a part of the whole powder for aluminum and blend soft aluminum powder into the aluminum alloy powder into which the whole amount of Zn is alloyed, in order to raise the compressibility of the raw material powder. For sufficiently achieving this purpose, the amount of used simple aluminum powder is necessary set to be 15 mass % of the whole raw material powder or more.
Regarding the aluminum alloy powder containing Zn, if it has a composition such that causes the production of Al—Zn liquid phase at a low temperature, Zn is likely to volatilize from this Al—Zn liquid phase. Therefore, it is preferable that the aluminum alloy powder has a composition with which the production of the Al—Zn liquid phase is caused at a temperature that is as high as possible, that is, only at a temperature of the final stage of the sintering step. Moreover, if using an aluminum alloy powder containing a large amount of Zn, this causes to relatively increase the amount of simple aluminum powder with the result that Zn dispersed in the sintered aluminum alloy matrix becomes likely not to be uniform. This causes the occurrence of fluctuation in the values of the obtained mechanical properties. In view of these items, it is preferable that the content of Zn in the aluminum alloy powder be 30 mass % or less. On the other hand, if the content of Zn in the aluminum alloy powder falls below 10 mass %, the difference in zinc concentration from the simple aluminum powder becomes small, with the result that Zn comes to have difficulty of being diffused and uniform dispersion is suppressed by contraries. Accordingly, it is preferable that the content of Zn in the aluminum alloy powder be in the range of from 10 to 30 mass %.
C. Forms of Mg and Cu
Use of the aluminum alloy powder having a composition with which the production of the Al—Zn liquid phase is caused only at a high temperature is preferable for preventing volatilization of zinc, but it is disadvantageous for uniform diffusion of the components. Cu and Mg are used together with Zn, for the purpose of causing the uniform diffusion of Zn into the above-described matrix. Cu and Mg, in the process wherein the temperature is elevated during sintering, cause the production of a Cu—Zn liquid phase or Mg—Zn liquid phase together with Zn powder or Zn in the aluminum alloy powder. These liquid phases are immediately solidified by their components being absorbed into the aluminum powder or aluminum alloy powder, and liquefaction and solidification are repeated so that uniformity of the components rapidly proceeds. Moreover, the liquid phase at this time gets solidified so immediately that no problems with the volatilization of Zn arise. The elements, Cu and Mg, each of which has the above-described action may be added in the form of simple metal powder, an alloy powder of the both elements, or an alloy powder with aluminum, and no hindrance occurs in any of the above cases. When the aluminum alloy powder containing Zn simultaneously contains Cu at the content of 10 mass % or less, the above-described effect becomes more enhanced. However, if the amount of Cu added into the aluminum alloy powder exceeds 10 mass % of the aluminum powder, the temperature at which Cu produces a liquid phase together with Zn shifts to the high-temperature side, and addition at more than 10 mass % is thus disadvantageous in terms of the uniform diffusion of the components.
D. Forms of Sn, Bi and In
If Sn, Bi or In is used as the sintering aid agent, they may be added in the form of simple metal powder, eutectic alloy powder or monotectic alloy powder that would cause the production of an eutectic liquid phase comprising those main components.
E. Powdered Material for Hard Particles
As the means for dispersing the hard particles into the aluminum alloy matrix, it is convenient to add a powdered material for the hard particles. If the powdered material reacts with the main component of the matrix, Al, it becomes difficult to control the amount and the range of particle size of the hard particles dispersed in the aluminum alloy matrix after the sintering. Therefore, it is preferred that the added powder, as the hard particles, is made of a material that does not react with aluminum.
For the hard particles as described above, silicon carbide, chromium boride, boron carbide and the like are preferable materials because they are extremely hard and does not react with aluminum. Since the aluminum alloy matrix is somewhat soft, the hard particles originated from a powder of the extremely hard material are embedded in the aluminum alloy matrix during the sliding operation, to suppress the wear of the sliding counterpart member, and, at the same time, they present plastic flow of the aluminum alloy matrix to contribute to improving wear resistance. Moreover, even when they once fall off the aluminum alloy matrix during the sliding operation, they are embedded again in the soft aluminum alloy matrix to repeatedly exhibit the effect of preventing plastic flow of the matrix.
(1)-3 Size of Powder
In order that the respective ingredient elements exhibit their roles uniformly in the matrix, it is necessary to uniformly diffuse those ingredient elements in the matrix. For this purpose, it is preferable that each of those ingredient elements be added in the form of fine powder whose particle size is as small as 74 microns (200 meshes) or less (i.e. 200 meshes minus sieve or the powder having a particle size that passes through a comb screen of 200 meshes), except for the simple aluminum powder. The simple metal powder or alloy powder, when the temperature is elevated during sintering, is melted to become a liquid phase, which wets the surface of the aluminum powder to eliminate the oxide layer and which is diffused into the aluminum matrix and simultaneously to accelerate the bond between the aluminum powder particles due to the diffusion. However, if the particle size of the simple metal powder or alloy powder exceeds 200 meshes, local segregation takes place, and uniform diffusion of the ingredient elements is obstructed.
However, if the aluminum powder is also made a fine powder, the flowability of the raw material powder becomes inferior. Therefore, it is suitable to use a powder for aluminum whose particle size is greater than that of the above-described respective ingredient element powder. Specifically, it is preferable to use a powder for aluminum whose particle size is 140 microns (100 meshes) or less (i.e. 100 meshes minus sieve or the powder whose particle is of a size having passed through a comb screen of 100 meshes). If exceeding the size of 100 meshes, each ingredient element has the difficulty of being diffused up to the center of the powder, and the component comes to get segregated. Therefore, such should be avoided.
Since the powdered material for the hard particles almost does not react to the matrix, it is to be dispersed in the aluminum alloy matrix and be left as it is added. Accordingly, the size of the raw powder used for hard particles can be determined as that of hard particles dispersed in the aluminum alloy matrix. The particle size of the hard particles dispersed in the aluminum alloy matrix is preferably 1 to 100 microns as the average. If the hard particles are smaller than 1 micron, they are easily flown with the matrix when the matrix flows plastically, and it is therefore difficult to prevent plastic flow of the matrix. Moreover, if the hard particles are larger than 100 microns, wear is easily caused on the sliding counterpart member during the sliding operation, depending on the sliding conditions. Therefore, in order to uniformly disperse in the aluminum alloy matrix the hard particles of a mean particle size in the above-mentioned range, it is preferred to use a powder of a material which does not react with aluminum, having a size of 113 microns (125 meshes) or less (i.e. 125 meshes minus sieve or the powder whose particle is of a size having passed through a comb screen of 125 meshes).
(2) Compacting Step
In this step, the raw material powder prepared from the above-described raw material powder blending step is filled into a die of a predetermined configuration, and the powder is then formed into a compact by compressing it under a compacting pressure of 200 MPa or more. As a result of this, a compact with a density ratio of 90% or more is obtained. If the compacting pressure falls below 200 MPa, the density of the compact becomes low, and, even after the compact passes through the subsequent sintering step and forging step, the pores of 2 vol % or more remain. This results in failure to impart high strength and elongation. Such insufficient compacting is not preferable also for the reason that dimensional change during sintering becomes large. The higher the compacting pressure is, the higher the density of the obtained compact becomes. Therefore, high compacting pressure is preferable. When the compacting pressure is 400 MPa or more, a compact whose density ratio is 95% or more is obtained and this is suitable. However, a compacting pressure exceeding 500 MPa easily causes adhesion of the aluminum powder to the die and it is therefore undesirable.
(3) Sintering Step
If a large amount of the relevant liquid phase mentioned above is produced during sintering, the amount of shrinkage of the sintered mass becomes large, with the result that the dimensional precision becomes inferior. Moreover, since zinc contained as an ingredient is an element having a low melting point and is therefore easy to volatilize in this sintering step, the amount of zinc that is dissolved in the base to make solid solution is reduced by the volatilization, resulting in failing to accomplish a desired value of strength and elongation. Simultaneously, zinc contaminates the sintering atmosphere and, in some cases, is deposited within the furnace, resulting in raising a problem with the working environment as well. To avoid inviting such bad effects, in a case of using a simple zinc powder, it is then recommended that elevation of the temperature up to the sintering temperature be performed at a high rate.
Namely, at the step of sintering the compact obtained in the above-described compacting step, it is recommended, in the course of temperature elevation from room temperature to the sintering temperature, that heating in the temperature range from at least 400 degrees C. being in the proximity of the melting point of zinc up to the sintering temperature is rapidly proceeded at the temperature-elevating rate of 10 degrees C./min or more to suppress the volatilization of the relevant ingredient elements. Moreover, sintering of the compact is developed by heating the compact at a sintering temperature of 580 to 610 degrees C. (with use of aluminum alloy powder containing zinc) or 590 to 610 degrees C. (with no Zn-containing aluminum alloy powder), for a sintering time of 10 minutes or more, so that, while the excessive decrease in the dimensional precision due to the generation of a liquid phase is being suppressed, uniform diffusion of the ingredient element is possibly achieved. If the temperature-elevating rate for elevating up to the sintering temperature is lower than 10 degrees C./min, the problem concerning volatilization of zinc becomes remarkable. If the sintering temperature exceeds 610 degrees C., the problems concerning volatilization of zinc and over-shrinkage due to the liquid phase become remarkable, and, in this case, the crystal grains also grow and become large, causing the decrease in the strength. It is necessary for uniform formation of solid solution with the respective ingredient elements in the Al base that the sintering temperature be settled to 580 degrees C. (using Zn-containing Al alloy powder), 590 degrees C. (using no Zn-containing Al alloy powder), or more, and that the sintering time length be made 10 minutes or more. If the sintering conditions fall below those ranges, diffusion of the respective ingredients into the Al base becomes insufficient, resulting in that the strength decreases.
By the above-described sintering, the respective ingredients are each kept in the state of their being dissolved in the matrix. The sintered compact is then cooled and the cooling rate had better be high although not particularly limited. In detail, if the cooling rate is low, in the high temperature range (450 degrees C. or more) in particular, the increase in size of the crystal grains proceeds. In addition, the component over-saturated in the course of cooling sometimes gets precipitated along the grain boundary, to cause the decrease in the strength and elongation. Also, that portion where the over-saturated component has been precipitated sometimes gets absorbed into the matrix by subjecting to a subsequent heat treatment (solution treatment), to make pores that cause the deterioration in the strength and elongation. Therefore, it is better to cool in the high temperature range at a rate that is as high as possible. Particularly, in the temperature range of 450 degrees C. or more, it is preferred that the sintered compact is cooled at a rate of −10 degrees C./min.
In regard to the sintering atmosphere, non-oxidizing one is suitable. Among various non-oxidizing gases, an atmosphere of nitrogen gas wherein the dew point is made —40 degrees C. or less is the most suitable. The dew point is an indicator that indicates the amount of water in the atmosphere of gas, and a large amount of water, that substantially means a large amount of oxygen, hinders the progress of the sintering operation since the Al is likely to have a bond to oxygen, to obstruct the densification of the mass. Since nitrogen gas is also inexpensive and safe comparing to other non-oxidizing gases, the nitrogen gas atmosphere that the dew point is specified as above is therefore preferable.
In accordance with the above sintering, the ingredient elements are uniformly dissolved in the Al matrix to make solid solution through liquid phase sintering, and a sintered compact such that the density ratio is 90% or more and the pores are closed pores is possibly obtained.
(4) Heat-Treating (T6 Treatment) Step
The heat-treating (T6 treatment in accordance with the regulation of JIS H 0001) step in the manufacturing method of the present invention comprises a solution treatment and an aging precipitation treatment. In the solution treatment, a precipitation phase in the Al base is uniformly dissolved in the Al base to form solid solution by heating at a temperature of from 460 to 490 degrees C., and the resulting mass is then water-quenched, thereby making an over-saturated solid solution. In the aging precipitation treatment, the resulting mass after the solution treatment is maintained at a temperature of 110 to 200 degrees C. for 2 to 28 hours to precipitate the over-saturated solid solution and form the precipitation phase dispersed in the Al base. If the temperature for the solution treatment is below 460 degrees C., the precipitated components does not uniformly form solid solution as a whole into the Al matrix. On the other hand, if that temperature exceeds 490 degrees C., although that effect almost does not change, a liquid phase is produced at a temperature exceeding 500 degrees C., to cause the generation of pores. In regard to the aging treatment, if the temperature is below 110 degrees C., or if the treatment time does not reach 2 hours, a sufficient amount of precipitated compound is not obtained, whereas, in a case where the temperature exceeds 200 degrees C., or a case where the treatment time exceeds 28 hours, the precipitated compound grows to become excessively large, resulting in the decrease in the strength. In regard to the length of time for the aging treatment, it is approximately 2 to 28 hours. The temperature and time length can suitably be adjusted, respectively, within the above-described ranges according to the property that is required. By subjecting to the above-described heat treatment, metallographic structure in which intermetallic compounds such as MgZn2 (η-phase), Al2Mg3Zn3 (T-phase), CuAl2 (θ-phase) are precipitated and dispersed in the aluminum alloy matrix is formed to achieve improvement of mechanical properties.
Regarding the wear-resistant sintered aluminum alloy with high strength, obtained through the above-described steps, it is so densified that the density ratio is 90% or more and it exhibits such an excellent property as a tensile strength of 450 MPa, as well as elongation and wear resistance that are equal to the conventional material. Moreover, it is possible to further improve the mechanical property, by an additional step for subjecting forging step between the sintering step and the heat treatment step.
(5) Forging Step Causing Plastic Flow Under Pressure
In this step, sinter forged aluminum parts exhibiting high tensile strength and high elongation can be obtained by subjecting the sintered mass obtained through the above-described steps before heat treatment and having a density ratio of 90%, to a cold forging step where it is forged at a room temperature at an upsetting ratio of 3 to 40%, or a hot forging step where it is forged at a temperature of 100 to 450 degrees C. at an upsetting ratio of 3 to 70%, to obtain a sinter forged aluminum part which has an increased density ratio of 98% or more. The resultant part has a high tensile strength and elongation.
In general, it is known to possibly increase the density through the execution of the forging treatment. However, in the case of porous material, simply increasing the density only results in that the pores get closed and no metallic bond is formed at the pore walls. As a result, that is followed by the occurrence during forging of cracks in the surface of the material or the remaining the pores as the defects within the product, failing to enhance the strength and elongation. Accordingly, in order to obtain a high level of strength and elongation, it is necessary not only to close the pores but also to form metallic bond there. To obtain this metallic bond, in general, the forging has been performed through two-divided sub-steps, one of which is a sub-step for performing densification of the relevant material and the other of which is a deforming sub-step for obtaining metallic bond by deforming the densified material.
In the present invention, for the process to obtain metallic bond, there is employed a technique of performing upsetting forging that comprises applying pressure, from the above and below, to the sintered porous material that has been obtained as above, to compress it in the direction of height for closing the pores, and also to deform the compressed material toward the space provided at the lateral side of the material for causing plastic flow of the material in the direction crossing to the direction in which the pressure is applied, thereby compulsively forming material bond of the original pore portions (i.e. the portion where the pore is closed although no metallic bond is made), while metallic bond is formed in these pore portions. Accordingly, the forging step of the present invention comprises a single operation into which the works of the two sub-steps that have been conventionally executed are merged. In connection with the upsetting forging, the upsetting ratio is determined as a ratio of the difference in the pressing direction between the dimensions before and after forging of the material relative to the dimension before forging of the material. Here, it is noted that the importance of the forging step of the present invention is to cause lateral plastic flow of the material under pressure. Therefore, if the above-described upsetting deformation is main work of the operation of the forging step, that is acceptable and no hindrance exists even when the operation of the forging step also locally or partially works as forward or backward extrusion on the material. Namely, the forging operation according to the present invention can include a technique wherein the material is locally extruded. Moreover, the processing operation that the area of material is reduced by means of a die, such as forging with forward or backward extrusion and the like, can also be included in the operation of the forging step, since the pressing in that operation works in radial directions and the direction in which the material is deformed is along the extruding direction or the one that intersects the pressing directions at right angles. Therefore, the above technique of working is also included in the scope of the present invention. Also, by performing the above forging operation for the compressing and plastically material flowing works described above, it is also possible to obtain, in addition to the above-described action, an action which makes fine the crystal grains that grew during sintering, as well as breaks the precipitate, whereby the strength and elongation are more enhanced.
In the case of cold forging, it is necessary to forge so that the upsetting ratio is 3 to 40%. If the upsetting ratio is less than 3%, deformation occurs only locally when the diameter after the forging is the same or larger in comparison with that before the forging, with the result that the amount of residual pores is increased to fail to enhance the strength and elongation. Also, in a case where forging is done with a die whose diameter is small such as forward extrusion forging, an upsetting ratio of 3% or more is necessary for the reason described above. Incidentally, if the upsetting ratio is 10% or more, the density ratio of the forged mass can easily be made to be 98% or more. Therefore, that setting is preferable. On the other hand, if the upsetting ratio exceeds 40%, cracking is likely to occur on the forged mass. When employing the cold forging operation, if upsetting forging is designed in such a manner that the terminal end portions of the material laterally extended during forging come to full contact with the inner wall of the die at the finish of the forging operation, the precision in dimension and shape of products increases and defects have difficulty remaining on the uppermost surface. Therefore, such way of upsetting forging is preferable.
In the case of hot forging, if heating the material (sintered mass) within a range of from 100 to 450 degrees C., preferably from 200 to 400 degrees C., forging at an upsetting ratio within a range of from 3 up to 70% is allowed. When the heating temperature for the material (sintered mass) is below 100 degrees C., almost no useful change is made in comparison with the case of cold forging. That is, the deformability of the material is still poor and it is therefore difficult to raise the upsetting ratio. In a case where the heating temperature of the material (sintered mass) is 200 degrees C. or more, the material becomes soft and the deformability increases. Accordingly, it is possible to decrease the forging pressure for performing hot forging at a desired value of upsetting ratio. Therefore, such temperature range is preferable. On the other hand, if the heating temperature exceeds 450 degrees C., the adhesion between the die and the material (sintered mass) remarkably occurs. Therefore, the upper limit needs to be set at 450 degrees C. at maximum, and preferably at 400 degrees C. However, even in the suitable temperature range described above, if the upsetting ratio exceeds 70%, forging cracks become likely to occur. Also in the case of hot forging, if upsetting forging is performed in such a manner that the terminal end portions of the material laterally extended during forging come to contact with the inner wall of the die at the finish of the forging operation, defects have difficulty remaining on the uppermost surface. Therefore, such way of upsetting forging is preferable.
The wear-resistant sinter forged aluminum alloy with high strength that have been obtained through the above-described process, as will be apparent from the following Examples, has a density ratio of 98% or more and is improved to have a tensile strength of 500 MPa or more and an elongation of 2% or more. Therefore, it exhibits high mechanical characteristics that cannot be expected from the one in the conventional art, as well as excellent wear resistance.
For each sample, the raw material powder blending step, the compacting step, the sintering step, the forging step, and the heat treatment step were sequentially performed to manufacture and evaluate five kinds of samples of sintered aluminum alloy having a overall composition shown in Table 2. Specifically, in the raw material powder blending step, aluminum powder having a particle size of minus sieve of 100 meshes screen; zinc powder, magnesium powder, copper powder, tin powder, bismuth powder, indium powder, lead-free solder powder containing 8 mass % Zn, 3 mass % Bi and the balance Sn, each of which had a particle size of minus sieve of 250 meshes screen respectively; silicon carbide powder, chromium boride powder and boron carbide powder, each of which had a particle size of minus sieve of 125 meshes screen, were prepared to provide a raw material powder by blending and mixing those powders together in accordance with the blending ratio shown in Table 1.
In the compacting step, adjusting the compacting pressure to 300 MPa, the raw material powder was formed into compacts of columnar shape having dimensions of +40 mm×28 mm for measuring mechanical property. In the sintering step, the compact was heated in an atmosphere of nitrogen gas, by elevating the heating temperature within a range of from 400° C. up to sintering temperature of 600 degrees C. at a temperature-elevating rate of 10 degrees C./min, and it was sintered by keeping it at the sintering temperature for 20 minutes. After that, the compact was cooled from the sintering temperature down to 450 degrees C. at a cooling rate of −20 degrees C./min. In the forging step, thus obtained sintered compact was heated at 400 degrees C. and put into a die of the same temperature to perform hot forging at an upsetting ratio of 40%. In the heat treatment step, the forged compact was heated at 470 degrees C. to perform the solution treatment, and it was then maintained at 130 degrees C. for 24 hours to perform aging precipitation treatment.
Moreover, as a conventional material, aluminum powder and aluminum-silicon alloy powder containing 20 mass % Si and the balance Al, each of which had a particle size of minus sieve of 100 meshes screen; nickel powder, copper-nickel alloy powder containing 4 mass % Ni and the balance Cu, and aluminum-magnesium alloy powder containing 50 mass % Mg and the balance aluminum, each of which had a particle size of minus sieve of 250 meshes screen, were prepared to provide a raw material powder by blending and mixing those powders together in accordance with the blending ratio shown in Table 1. In the compacting step, the compacting pressure was adjusted to 200 MPa, and, in the sintering step, the compact was heated in an atmosphere of nitrogen gas, by elevating the heating temperature within a range of from 400° C. up to sintering temperature of 550 degrees C. at a temperature-elevating rate of 10 degrees C./min, and the sintering temperature was kept for 20 minutes before cooling from the sintering temperature down to 450 degrees C. at a cooling rate of −20 degrees C./min. In the forging step, the heating temperatures of the sintered compact and the die were 450 degrees C., and the upsetting ratio was 40%. In the heat treatment step, the temperature for solution treatment was 470 degrees C., and the aging precipitation was performed at 130 degrees C. for 24 hours, to produce an alloy disclosed in the document of JPA No. H07-224341.
In the preparation of each sample, the density ratio was measured for each of the green compact after the compacting step, the sintered compact after the sintering step and the forged compact after the forging step. The results are shown in Table 3. For the evaluation of obtained samples of Nos. A01 to A34, five columnar pieces of φ40 mm×28 mm were processed into a tensile test piece and tensile test was conducted thereon to measure the tensile strength and elongation. The result is shown as an average value in Table 3. Moreover, each of another two columnar pieces were cut into an wearing test piece having a columnar shape of φ 7.98 mm×20 mm, sliding test was conducted thereon by a pin-on-disk wear resistance test machine at a sliding speed of 5 m/s for 30 min., using a counterpart member made of S45C heat-treated material under a constant load and supplying an engine oil thereto. If drastic change in the coefficient of dynamic friction was not observed during the sliding test, the test piece was substituted with another one and the load was raised by 5 MPa in every time. The load under which drastic increase in the coefficient of dynamic friction was observed was determined as a seizure pressure (critical bearing pressure). The results are shown together in Table 3.
Comparing samples of Nos. A01-A08 with one another, the effect of the hard particles with the addition amount is possibly researched. From the results, it is understood that the sample No. A01 containing no had particles exhibits high tensile strength and elongation but the seizure pressure is small, meaning a material having a low wear resistance. Even in such a material, the wear resistance can be improved by the hard particles at an amount of 0.1 mass % or more so that the seizure pressure rises to 30 MPa or more, while suppressing fall of the tensile strength to a small extent. In particular, addition at 1.0 mass % or more provides high wear resistance. On the other hand, the elongation tends to slightly decrease according as the amount of hard particles increases, but it is still possible to exhibit sufficient elongation, with an amount of 10 mass % or less of the hard particles. However, if the amount of hard particles exceeds 10 mass %, decrease in elongation becomes remarkable, and the wear amount of the counterpart member increases simultaneously. From the above, it is confirmed that, when the amount of hard particles is in a range of 0.1 to 10 mass %, high tensile strength and high elongation are exhibited, while the wear resistance is improved, resulting in provision of a wear-resistant sintered aluminum alloy exhibiting higher tensile strength than that of the wear-resistant sintered aluminum alloy of sample No. A33 which is a conventional aluminum-silicon alloy. It is also found that the effect of improving wear resistance is especially great when the amount of hard particles is in a range of 1.0 to 10 mass %.
Comparing sample of No. A06 with samples of Nos. A09-A12, the effect of zinc with the addition amount is possibly researched. From the results, it is understood that the amount of the intermetallic compound precipitated in the aluminum alloy matrix in the sample No. A09 containing 2.0 mass % zinc is so poor that the tensile strength is small, though the elongation is great. Moreover, its seizure pressure is of a low value, in spite of the hard particles added at 5.0 mass %. In contrast, at the zinc content of 3.0 mass % or more, although the elongation is reduced by the increase in the amount of intermetallic zinc compound precipitated in the aluminum alloy matrix, the tensile strength and the seizure pressure are increased by the effect of this increased amount of intermetallic compound and the effect of increased density of the sintered mass that is caused by the increased amount of Zn liquid phase and/or eutectic phase of Zn and other element. However, if the amount of precipitated intermetallic zinc compound is excessive, the strength is rather damaged, so that the tensile strength and the seizure pressure tend to be at peaks on sample No. A06 containing 5.5 mass % of Zn and then turn to decrease on sample No. All containing 10 mass % of Zn. Moreover, it has found, in sample No. A12 in which the Zn content exceeds 10 mass %, that the sintered compact was molten with excessive amount of zinc-containing liquid phase produced during sintering, resulting in canceling the subsequent forging step, heat treatment and tests. From the above, it can be confirmed that the zinc content of 3.0 to 10 mass % is effective in improvement of tensile strength and seizure pressure.
Comparing sample of No. A06 with samples of Nos. A13-A17, the effect of magnesium with the addition amount is possibly researched. From the results, it is understood that the amount of the intermetallic compound precipitated in the aluminum alloy matrix in the sample No. A13 containing 0.1 mass % Mg is so poor that the tensile strength is small, though the elongation is great. Moreover, its seizure pressure is of a low value, in spite of the hard particles added at 5.0 mass %. In contrast, at the magnesium content of 0.5 mass % or more, although the elongation is reduced by the increase in the amount of intermetallic compound precipitated in the aluminum alloy matrix, the tensile strength and the seizure pressure are increased by the effect of this increased amount of intermetallic compound and the effect of increased density of the sintered mass that is caused by the increased amount of Mg liquid phase and/or eutectic phase of Mg and other element. However, if the amount of precipitated intermetallic compound is excessive, the strength is rather damaged, so that the tensile strength and the seizure pressure tend to be at peaks on sample No. A06 containing 2.5 mass % of Mg and then turn to decrease on sample No. A16 containing 5 mass % of Mg. Moreover, it has found, in sample No. A17 in which the Mg content exceeds 5.0 mass %, that the sintered compact was deformed with excessive amount of zinc-containing liquid phase produced during sintering, resulting in canceling the subsequent forging step, heat treatment and tests. From the above, it can be confirmed that the magnesium content of 0.5 to 5.0 mass % is effective in improvement of tensile strength and seizure pressure.
Comparing sample of No. A06 with Nos. A18-A22, the effect of copper with the addition amount is possibly researched. From the results, it is understood that the amount of the intermetallic compound precipitated in the aluminum alloy matrix in the sample No. A18 containing 0.1 mass % Cu is so poor that the tensile strength is small, though the elongation is great. Moreover, its seizure pressure is of a low value, in spite of the hard particles added at 5.0 mass %. In contrast, at the copper content of 0.5 mass % or more, although the elongation is reduced by the increase in the amount of intermetallic compound precipitated in the aluminum alloy matrix, the tensile strength and the seizure pressure are increased by the effect of this increased amount of intermetallic compound and the effect of increased density of the sintered mass that is caused by the increased amount of Cu liquid phase and/or eutectic phase of Cu and other element. However, if the amount of precipitated intermetallic compound is excessive, the strength is rather damaged, so that the tensile strength and the seizure pressure tend to be at peaks on sample No. A06 containing 1.5 mass % of Cu and then turn to decrease on sample No. A21 containing 5.0 mass % of Cu. Moreover, it has found, in sample No. A22 in which the Cu content exceeds 5.0 mass %, that the sintered compact was deformed with excessive amount of zinc-containing liquid phase produced during sintering, resulting in canceling the subsequent forging step, heat treatment and tests. From the above, it can be confirmed that the copper content of 0.5 to 5.0 mass % is effective in improvement of tensile strength and seizure pressure.
Comparing sample of No. A06 with Nos. A23 and A24 in Tables 1-3, the effect of the hard particles with the kind of them is possibly researched. From the results, it is understood that sufficient wear resistance (seizure pressure) is possibly achieved even if the kind of hard particles is changed from boron carbide to silicon carbide or chromium boride. It has been found that, especially when chromium boride is used, it is possible to provide an excellent wear-resistant sintered aluminum alloy which exhibits not only a higher tensile strength than that of the wear-resistant sintered aluminum alloy (sample No. A33) of the conventional aluminum-silicon type, but also an equal seizure pressure.
Comparing sample of No. A06 with Nos. A25-A29 in Tables 1-3, the effect of the low-melting-point metal powder with the addition amount is possibly researched. From the results, it is understood that sufficient tensile strength, elongation and seizure pressure are possibly achieved even if the low-melting-point metal powder is not added, and that, in particular, the tensile strength and the elongation are higher than those of the wear-resistant sintered aluminum alloy (sample No. A33) of the conventional aluminum-silicon type. Moreover, it is also understood that these characteristics are improved by adding 0.01 to 0.5 mass % of low-melting-point metal powder to the wear-resistant sintered aluminum alloy of the present invention. However, if its addition exceeds 0.5 mass %, the low-melting-point metal precipitates in the grain boundary of the aluminum alloy matrix and the above characteristics rather deteriorate. From the above, it can be confirmed that, though no addition of the low-melting-point metal powder is allowable, its addition at 0.01 to 0.5 mass % is effective to improve the tensile strength elongation and seizure pressure.
Comparing sample of No. A06 with Nos. A30-A32 in Tables 1-3, the effect of the low-melting-point metal powder with the kind of it is possibly researched. From the results, it is understood that, even if the kind of the low-melting-point metal powder is changed from tin to bismuth, indium or eutectic compounds thereof (lead-free solder), the same improving effect as that with tin can be obtained in tensile strength, elongation and seizure pressure.
In this example, using the raw material powder prepared in Example 1 and the same blending ratio for sample No. A03: 5.5 mass % zinc powder; 2.5 mass % magnesium powder; 1.5 mass % copper powder; 5 mass % boron carbide powder; 0.1 mass % tin powder; and the balance aluminum powder, sintered aluminum alloy samples were manufactured by performing the same operation of Example 1, excepting that the compacting pressure, the sintering conditions (temperature-elevating rate in the range from 400 degrees C. to the sintering temperature, the sintering temperature and time), and the forging conditions (the heating temperatures of the sintered compact and the forging die, the upsetting ratio) were changed to those shown in Table 4. Regarding each of these samples, the same estimation in Example 1 was executed. The results are shown in Table 5.
Comparing sample of No. A06 with Nos. A34-A37 in Tables 4 and 5, the effect of the compacting pressure is possibly researched. From the results, it is understood that a compacting pressure in a rage of 200 to 400 MPa provides a green compact sample having a high density ratio, which results in a sintered aluminum alloy exhibiting high tensile strength and high elongation through the process of sintering-forging-heat treatment. In contrast, the green compact of sample No. A34 in which the compacting pressure is less than 200 MPa has a low density, and its sintered mass thus caused deformation by large shrinkage due to generation of liquid phase, resulting in canceling the subsequent forging step, heat treatment and tests. On the other hand, in the green compact of sample No. A37 in which the compacting pressure exceeds 400 MPa, adhesion (die galling) of the sintered compact to the die has occurred when taking the compact out of the die, also resulting in canceling the subsequent forging step, heat treatment and tests. From the above, it is confirmed that the compacting is necessarily conducted at a compacting pressure of 200 to 400 MPa.
Comparing sample of No. A06 with Nos. A38-A40 in Tables 4 and 5, the effect of the temperature-elevating rate in the range of 400 degrees C. up to the sintering temperature is possibly researched. From the results, it is understood that, in regard to the sample of No. A38 in which the temperature-elevating rate is less than 10 degrees C./min, the Zn component volatilizes from the compact during sintering and the quantity of precipitation phase decreases, resulting in deterioration in tensile strength, elongation and seizure pressure. On the other hand, regarding the samples in which the temperature-elevating rate is 10 degrees C./min or more, it is seen that each of them exhibits a high level of tensile strength, elongation and seizure pressure. From the above, it is confirmed that the temperature-elevating rate in the range of 400 degrees C. to the sintering temperature is necessarily 10 degrees C./min.
Comparing sample of No. A06 with Nos. A41-A44 in Tables 4 and 5, the effect of the sintering temperature is possibly researched. From the results, it is understood that, regarding each of the samples in which the sintering temperature is within a range of from 590 to 610 degrees C., the sample exhibits high levels of tensile strength, elongation and seizure pressure. In contrast, regarding the sample of No. A41 in which the sintering temperature is lower than 590 degrees C., both the tensile strength and the elongation are decreased. The reason for this is considered that the ingredient element added in the form of a simple metal powder is not completely dissolved in the Al base and forms solid solution which is locally segregated to remain, with the result that the relevant sample has a small value of mechanical property. Conversely, in the sample of No. A44 in which the sintering temperature is over 610 degrees C., melting of the sintered compact occurs due to the liquid phase generating in excess. The succeeding test for evaluation has been therefore canceled. From the above, it is confirmed that the sintering temperature is necessarily in a range of 590 to 610 degrees C.
Comparing sample of No. A06 with Nos. A45-A48 in Tables 4 and 5, the effect of the sintering time is possibly researched. From the results, it is understood that, in the sample of No. A45 in which the sintering time is less than 10 minutes, it is low in terms of all of tensile strength, elongation and seizure pressure. The reason of this is considered that, when the sintering time is short, the ingredient is not sufficiently dissolved in the Al base and forms solid solution which is locally segregated to remain, with the result that the relevant sample has a small value of mechanical property. On the other hand, in the samples in which the sintering time is 10 minutes or more, the ingredient is uniformly dissolved in the Al base to form solid solution. Therefore, the relevant samples exhibit high levels of tensile strength, elongation and seizure pressure. However, even if the sintering time exceeds 30 minutes, these properties are not changed very much. Therefore, the setting of the sintering time being 30 minutes or less will be sufficient.
Comparing samples of Nos. A49-A66 in Tables 4 and 5 with one another, the effect of performing the forging or not, and the effects of forging temperature and upsetting ratio are possibly researched.
First, by comparing the sample of No. A49 that is the wear-resistant sintered aluminum alloy according to the present invention, with the sample of No. A66 that is a wear-resistant sintered aluminum alloy of the conventional aluminum-silicon type, in both of which the forging is not carried out, each of the elongation and the seizure pressure is of the same level in both samples. However, in regard to the tensile strength, it is confirmed that the sample of No. A49 of the present invention is excellent to exhibit a higher value than that of the conventional one.
Moreover, comparing with the sample of No. A49 in which the forging has not been carried out, the samples of Nos. A50-65 (excepting ones that the test for evaluation was canceled for improperness) in which the forging has been performed are improved in all of tensile strength, elongation and seizure pressure. Therefore, the effect of adding the forging step has been confirmed.
Second, the forging conditions are searched. Comparing the samples of Nos. A49-A54 with one another, if the upsetting ratio is in a range from 3 to 40% in the case of cold forging, it is found that the improving effect can be seen in all of tensile strength, elongation and seizure pressure. In contrast, if the upsetting ratio exceeds 40% such as in the sample of No. A54, cracks occur in the sample due to forging. The test for evaluation of the sample was therefore canceled.
Comparing the sample of No. A35 (cold forging) with Nos. A06 and A55-A60, it is also understood that, in the case of hot forging with changing the temperatures of the sintered mass and the forging die, the tensile strength is possibly improved, while the elongation is largely improved. This is attributable to the fact that, although in the case of cold forging hair cracks very slightly remain within the sample, followed by decrease in the elongation, carrying out hot forging of the material with the heating temperature being set to 100 degrees C. or more makes the hair cracks removed. On the other hand, when the forging temperature exceeds 450 degrees C., adhesion (die galling) of the sintered compact to the die occurs. Therefore, the test in such a case has been cancelled.
Third, comparing the sample of No. A06 and the samples of Nos. A61-A65, it is understood, even when the upsetting ratio within a wide range of 3 to 70% is employed, the effect of improvement can be seen in each of tensile strength, elongation and seizure pressure. On the other hand, when the upsetting ratio exceeds 70% such as the sample of No. A65, forging causes cracks in the sample. Therefore, the test in such a case has been cancelled.
As described above, it is confirmed that the effect of improvement is obtained in tensile strength, elongation and seizure pressure, by adding, after the sintering step, either of the cold forging step where the sintered compact is forged at a room temperature and an upsetting ratio of 3 to 40%, and the hot forging step where the sintered compact is forged at a temperature of 100 to 450 degrees C. and an upsetting ratio of 3 to 70%.
In this example, comparison was made between the case where zinc is incorporated in the form of aluminum alloy powder (No. B01) and the case where it is in the form simple zinc powder (No. B02). Specifically, in the raw material powder blending step for No. B01, aluminum powder having a particle size of minus sieve of 100 meshes screen; aluminum alloy powder containing 12 mass % Zn; boron carbide powder having a particle size of minus sieve of 125 meshes screen as a powder for the hard particles; zinc powder, magnesium powder, copper powder and tin powder, each of which had a particle size of minus sieve of 250 meshes screen respectively, were prepared to provide a raw material powder having a overall composition of Zn: 5.5%, Mg: 2.5%, Cu: 1.5%, Sn: 0.1%, hard particles (boron carbide): 5.0% and the balance Al and inevitable impurities, by blending and mixing those powders together in accordance with the blending ratio shown in Table 6.
In the compacting step, adjusting the compacting pressure to 300 MPa, the raw material powder was formed into compacts of columnar shape having dimensions of φ40 mm×28 mm. In the sintering step, the compact was heated in an atmosphere of nitrogen gas, by elevating the heating temperature within a range of from 400° C. up to sintering temperature of 600 degrees C. at a temperature-elevating rate of 10 degrees C./min, and it was sintered by keeping it at the sintering temperature for 20 minutes. After that, the compact was cooled from the sintering temperature down to 450 degrees C. at a cooling rate of −20 degrees C./min. In the forging step, thus obtained sintered compact was heated at 400 degrees C. and put into a die of the same temperature to perform hot forging at an upsetting ratio of 40%. In the heat treatment step, the forged compact was heated at 470 degrees C. to perform the solution treatment, and it was then maintained at 130 degrees C. for 24 hours to perform aging precipitation treatment.
For the evaluation of each of the obtained samples of Nos. B01 and B02, five columnar pieces of φ40 mm×28 mm were processed into a tensile test piece and tensile test was conducted thereon to measure the tensile strength and elongation. The result is shown as an average value and a value 3a in Table 7. Moreover, each of another two columnar pieces were cut into an wearing test piece having a columnar shape of φ7.98 mm×20 mm, and sliding test was conducted thereon by a pin-on-disk wear resistance test machine at a sliding speed of 5 m/s for 30 min., using a counterpart member made of S45C heat-treated material under a constant load and supplying an engine oil thereto. If drastic change in the coefficient of dynamic friction was not observed during the sliding test, the test piece was substituted with another one and the load was raised by 5 MPa in every time. The load under which drastic increase in the coefficient of dynamic friction was observed was determined as a seizure pressure (critical bearing pressure). The results are shown together in Table 7. Moreover, in the preparation of each sample, the density ratio was measured for each of the green compact after the compacting step, the sintered compact after the sintering step and the forged compact after the forging step. The results are shown in Table 7. It is noted that the conditions in the compacting to aging precipitation treatment for sample No. B02 are the same as those for sample No. A06.
From the results in Tables 6 and 7, it is confirmed that the tensile strength becomes slightly higher and that the fluctuation in terms of the tensile strength can be particularly suppressed to a smaller range of values in a case where Zn is added in the form of the alloyed powder with Al (the sample No. B01), than in a case where Zn is added in the form of simple component powder (the sample No. B02). Moreover, the elongation is also improved and its fluctuation is suppressed to a small range. This is considered as the effect of that volatilization of zinc is prevented by adding in the alloy form the zinc component which is easily volatilized and the zinc content in the sample is thus not fluctuated. In contrast, an equivalent value is obtained for the seizure pressure. In this example, it is confirmed that improvement and suppressing of fluctuation in tensile strength and elongation can be achieved, without leading to decrease in seizure pressure, by adding zinc in the alloy form.
In the raw material powder blending step, aluminum powder having a particle size of minus sieve of 100 meshes screen; aluminum alloy powder of the composition shown in Table 8; magnesium powder, copper powder, and tin powder as a low-melting-point metal powder, each of which had a particle size of minus sieve of 250 meshes screen respectively; and boron carbide powder having a particle size of minus sieve of 125 meshes screen as a powder for the hard particles, were prepared to provide a raw material powder, by blending and mixing those powders together in accordance with the blending ratio shown in Table 8. Using each of these raw material powders, the compacting step, sintering step, forging step and heat treatment step were executed under the same conditions as in Example 3 to prepare each of samples having a overall composition shown in Table 9.
In the preparation of each sample, the density ratio was measured for each of the green compact after the compacting step, the sintered compact after the sintering step and the forged compact after the forging step, and measurement of tensile strength, elongation and seizure pressure (critical bearing pressure) was also conducted, the results of which are shown in Table 10.
Comparing the samples of Nos. B01 and B03-B08 in Tables 8 to 10, the effect of the amount of aluminum powder added is searched. In the samples of Nos. B03, and B04 wherein the amount of aluminum powder added is less than 15 mass %, as a result of the fact that the amount of Zn in the overall composition of the raw material powder becomes excessively large to such an extent as it exceeds 10 mass %, the sintered compact is largely deformed due to the liquid phase occurring from inside the aluminum alloy powder. The subsequent steps have therefore been canceled. On the other hand, when the amount of aluminum powder added is over 15 mass %, it becomes possible to sinter without occurring deformation of the sintered mass and the sample exhibits high levels of tensile strength, elongation and seizure pressure. From the above-mentioned results, it is confirmed that, in a case where Zn is wholly added in the form of aluminum alloy powder, it is necessary to simultaneously use the aluminum powder of 15 mass % or more. If the amount of aluminum powder added exceeds 15 mass %, the relevant sample also tends to exhibit enhanced values of tensile strength and elongation as the amount of aluminum powder increases. However, if the amount of Zn in the overall composition of the raw material powder exceeds 5.5 mass % (of sample No. B01), the tensile strength tends to fall on the contrary. In the sample of No. B08 wherein the amount of Zn in the overall composition of the raw material powder is lower than 3 mass %, with the result that the amount of zinc is poor, the decrease in the tensile strength and seizure pressure of the relevant sample is seen.
By comparing the samples of Nos. B01 and B09 to B14 in Tables 8 to 10, the effect of the content of Zn in the aluminum alloy powder is searched. In these comparisons, the amount of Zn in the overall composition of the raw material powder in each sample has been adjusted to a fixed value. From the results of these samples, it is found that, in the sample of No. B09, wherein the content of Zn in the aluminum alloy powder is less than 10 mass %, the product exhibits a high value of tensile strength whereas the elongation value thereof is small or 0.7%. On the other hand, in a case where the content of Zn in the aluminum alloy powder is 10 mass % or more, it is found that not only does the relevant sample exhibit a high tensile strength but is the value of elongation also enhanced. However, when the content of Zn in the aluminum alloy powder exceeds 30 mass %, both the decrease in the tensile strength and the decrease in the elongation are seen to occur (sample No. B14). Moreover, in regard to the seizure pressure, a preferable value is obtained when the content of Zn in the aluminum alloy powder is in a range of 10 to 30 mass %, but the decrease in the seizure pressure is seen when the content of Zn in the aluminum alloy powder exceeds 30 mass %. Accordingly, it is confirmed that, when the amount of Zn in the aluminum alloy powder is in the range of from 10 to 30 mass %, the relevant sample exhibits high values of tensile strength, elongation and seizure pressure.
In the optimum range confirmed as above of Zn content in the aluminum alloy powder, the lower limit value of Zn in the overall composition of the raw material powder, and the upper limit value thereof, can be searched, respectively, by the sample No. B15 and the sample No.B16. As a result, it is confirmed that, when the amount of Zn is in the range of from 3 to 10 mass % in the overall composition of the raw material powder, the relevant samples exhibit a high tensile strength, high elongation together and high seizure pressure with the above-described effect.
Example 5 is an embodiment wherein examination has been performed of the amounts of Mg and Cu added and the forms in which Mg and Cu were added. In this example, together with the aluminum powder, aluminum alloy powder, magnesium powder, copper powder, tin powder and boron carbide powder used in Example 3, mixed together were the aluminum alloy powders each having a composition shown in Table 11 and a particle size of 100 meshes minus sieve and the aluminum-magnesium alloy powder wherein the content of Mg was 50 mass %, the balance being Al and inevitable impurities and the particle size was 250 meshes minus sieve. The blending proportion is shown in Table 11, and the raw material powders each having an overall composition shown in Table 12 were prepared. Using these raw material powders, there were executed the compacting step, sintering step, forging step, heat-treating step and test piece processing step, under the same conditions as those in Example 3. Regarding the samples of Nos. B-17 to B-32 that were obtained above, the density ratios in each step as well as the mechanical properties, namely, tensile strength, elongation and seizure pressure, were measured, the results being shown in Table 13 together with the measured result (average value) of the sample No. B01 in Example 3.
By comparing the samples of Nos. B01, B17-B19, B21 and B22 in Tables 11 to 13, the effect of the amount of Mg powder that is added in the form of a simple metal powder is searched. From the results, it is found that, in the case of the Mg being not added whatsoever (the sample No. B17), where the liquid phase that Mg would otherwise participate in does not occur, the tensile strength, the elongation and the seizure pressure are reduced. In contrast, in a case where Mg is added in the form of a simple metal powder, the tensile strength, the elongation and the seizure pressure are enhanced when the amount of Mg is 0.5 mass % or more. However, in the case of sample No. B22 wherein the amount of Mg exceeds 5 mass %, the amount of liquid phase occurring becomes excessively large, with the result that the sintered compact is deformed. From these items, it is confirmed that, regarding the amount of Mg in the overall composition of the raw material powder, there is the effect of enhancing all of the tensile strength, the elongation and the seizure pressure when the amount of Mg is in the range of from 0.5 to 5 mass %.
Sample No. B20 is an example wherein Mg is added in the form of aluminum-magnesium alloy powder. Comparing it with the sample of No. B01, it is found that, if the amount of Mg is equal in the overall composition of the raw powder, the equivalent values of tensile strength, elongation and seizure pressure are obtained even when Mg is added in the form of aluminum-magnesium alloy powder.
By comparing the samples of Nos. B01 and B23-B27 in Tables 11 to 13, the effect of the amount of Cu powder that is added in the form of a simple metal powder is searched. From the results, it is found that, in the case of the Cu being not added whatsoever (sample No. B23), wherein the liquid phase that Cu would otherwise participate in does not occur, both of the tensile strength and the seizure pressure has a low value. In contrast, in a case where Cu is added in the form of a simple metal powder, the tensile strength and the seizure pressure are enhanced when the amount of Cu is 0.5 mass % or more. However, in the case of sample No. B27 wherein the amount of Cu exceeds 5 mass %, the amount of liquid phase occurring becomes excessively large, with the result that the sintered compact is deformed. On the other hand, regarding the elongation, as the amount of Cu increases, the elongation tends to decrease, but it is possibly held to be 1.0% or more at the amount of Cu in a range of up to 5 mass %. From these items, it is confirmed that, regarding the amount of Cu in the overall composition of the raw material powder, there is the effect of enhancing the tensile strength and the seizure pressure when the amount of Cu is in the range of from 0.5 to 5 mass %, and the amount of Cu in this range is preferable, because a sufficient value of elongation is obtainable.
By comparing the samples of Nos. B28-B32 in Tables 11 to 13, the effect of the amount of Cu in a case where Cu is added in the form of an aluminum alloy powder containing Zn therein is searched. In this case, as in the case where Cu is added in the form of a simple metal powder, the enhancement in the tensile strength and the seizure pressure is seen in comparison with the product wherein Cu is not added at all (the sample No. B23). However, regarding the amount of Cu in the overall composition of the raw material powder, even when it falls within the range of from 0.5 to 5 mass % that has been confirmed above, it is seen that, if the amount of Cu in the aluminum alloy powder exceeds 10 mass %, the tensile strength and elongation become contrarily decreased. From this result, it is further confirmed that, in a case where Cu is added in a form wherein it is alloyed into the aluminum alloy powder containing Zn therein, the upper limit of Cu in the alloy needed to be 10 mass %.
Example 6 is an embodiment wherein examination has been performed of the amount of hard particles and the kind thereof. Together with the aluminum powder, aluminum alloy powder, magnesium powder, copper powder, tin powder and boron carbide powder of the Example 3, used were the silicon carbide powder and the chromium boride powder each having a particle size of 125 meshes minus sieve. These powders were mixed together in the proportion for blending shown in Table 14, to prepare raw material powders each having an overall composition shown in Table 15. Using these raw material powders, there were executed the compacting step, sintering step, forging step, heat-treating step and test piece processing step, under the same conditions as those in Example 3, to obtain the products of sample Nos. B33-B42. Regarding the obtained samples, measurement of the density ratios in each step as well as the tensile strength, elongation and seizure pressure was carried out, the results being shown in Table 16 together with the measured result (average value) of the sample No.B01 in Example 3.
Moreover, as a conventional material, aluminum powder and aluminum-silicon alloy powder containing 20 mass % Si and the balance Al, each of which had a particle size of minus sieve of 100 meshes screen; nickel powder, copper-nickel alloy powder containing 4 mass % Ni and the balance Cu, and aluminum-magnesium alloy powder containing 50 mass % Mg and the balance aluminum, each of which had a particle size of minus sieve of 250 meshes screen, were prepared to provide a raw material powder by blending and mixing those powders together in accordance with the blending ratio shown in Table 6. In the compacting step, the compacting pressure was adjusted to 200 MPa, and, in the sintering step, the compact was heated in an atmosphere of nitrogen gas, by elevating the heating temperature within a range of from 400° C. up to sintering temperature of 550 degrees C. at a temperature-elevating rate of 10 degrees C./min, and the sintering temperature was kept for 60 minutes before cooling from the sintering temperature down to 450 degrees C. at a cooling rate of −20 degrees C./min. In the forging step, the heating temperatures of the sintered compact and the die were 450 degrees C., and the upsetting ratio was 40%. In the heat treatment step, the temperature for solution treatment was 470 degrees C., and the aging precipitation was performed at 130 degrees C. for 24 hours, to produce an alloy disclosed in the document of JPA No. H07-224341. Also for this sample (No. B43), measurement of the density ratio after each step as well as mechanical properties, namely, the tensile strength, elongation and seizure pressure was carried out. These results are also shown in Table 16.
Comparing the samples of Nos.B01 and B33-B40 in Tables 15 to 16, the effect of the amount of the hard particles is searched. From the results, it is understood that the sample No. B33 containing no had particles exhibits high tensile strength and elongation but the seizure pressure is small, meaning a material having a low wear resistance. Even in such a material, the wear resistance can be improved by the hard particles at an amount of 0.1 mass % or more so that the seizure pressure is raised, while suppressing fall of the tensile strength to a small extent. In particular, addition at 1.0 mass % or more provides high wear resistance. On the other hand, the elongation tends to slightly decrease according as the amount of hard particles increases, but it is still possible to exhibit sufficient elongation of 1% or more, with an amount of 10 mass % or less of the hard particles. However, if the amount of hard particles exceeds 10 mass % (sample No. B40), decrease in elongation becomes remarkable to fall below 1%, and it has bee simultaneously observed that the wear amount of the counterpart member has increased. From the above, it is confirmed that, when the amount of hard particles is in a range of 0.1 to 10 mass %, high tensile strength and high elongation are exhibited, while the wear resistance is improved, resulting in provision of a wear-resistant sintered aluminum alloy exhibiting higher tensile strength than that of the wear-resistant sintered aluminum alloy of sample No. B43 which is a conventional aluminum-silicon alloy. It is also found that the effect of improving wear resistance is especially great when the amount of hard particles is in a range of 1.0 to 10 mass %.
By comparing samples of Nos. B01, B41 and B42 in Tables 14 to 16, the effect of the hard particles with the kind of them is possibly researched. From the results, it is understood that sufficient wear resistance (seizure pressure) is possibly achieved even if the kind of hard particles is changed from boron carbide to silicon carbide or chromium boride. It has been found that, especially when chromium boride is used, it is possible to provide an excellent wear-resistant sintered aluminum alloy which exhibits not only a higher tensile strength than that of the wear-resistant sintered aluminum alloy (sample No. B43) of the conventional aluminum-silicon type, but also an equivalent value of seizure pressure.
Example 7 is an embodiment wherein examination has been performed of the amounts of sintering aid powder and the kind thereof. Together with the aluminum powder, aluminum alloy powder, magnesium powder, copper powder, boron carbide powder and tin powder of the Example 3, used were the bismuth powder, indium powder and the lead-free solder powder each having a particle size of 250 meshes minus sieve, and the lead-free solder powder had a composition wherein the content of Zn was 8 mass % and the amount of Bi was 3 mass %, the balance being Sn and inevitable impurities. These powders were mixed together in the proportion for blending shown in Table 17, to prepare raw material powders each having an overall composition shown in Table 17. Using these raw material powders, there were executed the compacting step, sintering step, forging step, heat-treating step and test piece processing step, under the same conditions as those in Example 3, to obtain the products of sample Nos. B44 to B51. Regarding the obtained samples, measurement of the density ratios in each step as well as the tensile strength, elongation and seizure pressure was carried out, the results being shown in Table 19 together with the measured result (average value) of the sample No.B01 in Example 3.
Here, comparing the samples of Nos. B01 and B44-B48 in Tables 17 to 19, the effect of the amount of the low-melting-point metal powder. Comparing with the product (sample No. B44) wherein no low-melting-point metal is added, it is found that, when the low-melting-point metal is added, the tensile strength, elongation and seizure pressure are improved. It is also found that, regarding that amount of addition, the effect of that is seen when that is in the range of from 0.01 to 0.5 mass %; and the effect is the highest when the adding amount thereof is in the range of from 0.05 to 0.1 mass %. However, if the adding amount thereof exceeds 0.5 mass %, the decrease in the elongation is outstanding, simultaneously accompanied by the decrease in the seizure pressure. Accordingly, it is confirmed that, regarding the addition of the low-melting-point metal powder, the effect of enhancing the mechanical properties is brought about when that addition is in the range of from 0.01 to 0.5 mass %.
Also, comparing the samples of Nos. B01 and B49-B51 in Tables 17 to 19, wherein the kind of the low-melting-point metal powder is changed, the effect of the kind of the low-melting-point metal powder is searched. From the results of them, it is confirmed that the same effect as described above is obtained even when the bismuth powder, indium powder or lead-free solder powder is used in place of tin powder.
Example 8 is an embodiment wherein examination is performed on a case where the compacting pressure is changed as a compacting condition, or one of the sintering temperature and sintering time is changed as a sintering condition.
Using the raw material powder which was prepared by using aluminum powder, aluminum alloy powder, magnesium powder, copper powder and tin powder and by adjusting to the same ingredient composition as that in Example 3, there were executed the compacting step and sintering step with the use of the compacting pressure, sintering temperature and sintering time shown in Table 20. Then, under the same conditions as those in Example 3, the forging step, heat-treating step and test piece processing step were performed. Regarding each of the obtained products, measurement of the density ratio in each step and the tensile strength, elongation and seizure pressure was carried out. The results are shown in Table 21 together with the result (average value) of sample No. B01 in Example 3.
From the results of samples of Nos. B01 and B52-B55 in Tables 20 and 21, it is found that, when the compacting pressure is in the range of from 200 to 400 MPa, a compacted compact sample in which the density ratio thereof is 80% or more, and that, by passing through the sintering-forging-heat treating steps, the product of the relevant sample exhibits a high level of tensile strength and high values of elongation and seizure pressure. Moreover, in the sample of No. B52 wherein the compacting pressure is below 200 MPa, the amount of shrinkage due to the occurrence of the liquid phase is large, because the density of the green compact is low. This has caused to lose the shape. As a result of this, the succeeding forging and heat-treating steps have been canceled and the relevant test has also been stopped. On the other hand, if the compacting pressure exceeds 400 MPa (in sample No. B55), die galling occurs, whereby the succeeding sintering step and the steps thereafter have been canceled and the test has been stopped.
Moreover, comparing the samples of Nos. B01 and B56-B59 in Tables 20 and 21, the effect of the sintering temperature is searched. From those results, it is found that the samples of Nos. B01, B67 and B58 wherein the sintering temperature is in the range of from 580 to 610 degrees C. exhibit a high level of tensile strength and a high value of elongation. On the other hand, in the sample of No. B56 wherein the sintering temperature is lower than 580 degrees C., both of the tensile strength and elongation are deteriorated. This is considered, because the ingredient element is not completely be dissolved in the Al base to form solid solution and it is locally segregated to remain, with the result that the mechanical properties deteriorate to a low value. Contrary to the above, in the sample of No. B59 wherein the sintering temperature is higher than 610 degrees C., the sintered compact is deformed with melting, because the amount of liquid phase excessively occurred. The succeeding test has been therefore canceled.
Comparing the samples of Nos. B01 and B60-B63 in Tables 20 and 21, the effect of the sintering time is searched. From the results of those, it is found that, in the sample of No. B60 wherein the sintering time is shorter than 10 minutes, both of the tensile strength and elongation are deteriorated. This is considered because the ingredient element is not sufficiently dissolved in the Al base to form solid solution and it is locally segregated to remain, with the result that the mechanical properties come to a low value. Opposite to the above, in the samples of Nos. B01 and B61-B63 wherein the length of sintering time is longer than 10 minutes, the ingredient is evenly dissolved in the Al base to form solid solution, whereby the relevant product exhibits a high level of mechanical property that, while the tensile strength is 500 MPa or more, the elongation exceeds 3%. Here, it is noted that, even when the sintering time exceeds 30 minutes, the mechanical property that the product exhibits has no change. Therefore, a sintering time of 30 min or less can be regarded as being sufficient.
In Example 9, the operation of Example 3 was repeated under the same conditions for sample production as those in Example 3, excepting that the forging conditions were changed to those shown in Table 22, to prepare product samples of Nos. B-53 to B-69, using the aluminum powder, aluminum alloy powder, magnesium powder, copper powder and tin powder used for the Sample No. B01 in Example 3 and preparing the raw material powders that were adjusted to the same ingredient composition as that in Example 3. Regarding each of these samples, the density ratio after executing each step as well as the tensile strength, elongation and seizure pressure was measured, the results being shown in Table 23 together with the measured results concerning the sample No. B01 in Example 3. Here, in Table 22, regarding the field “Forging Temperature”, the term “R.T. (Room Temperature)” designates the case of cold forging, and, in the case of hot forging, the heating temperature for a sintered compact sample as a material to be forged is shown. The sample of No. B64 is prepared for comparison with a specimen of a conventional material that is similar to the material of Japanese Laid-Open Patent Application of Publication No. JPA H04-365832 with no forging.
Here, comparing the samples of Nos. B64-B69 in Tables 22 and 23, the effect of the upsetting ratio that is brought about when cold forging is done at room temperature is searched. From those results, it is found that, in the case of cold forging, the sample has high levels of tensile strength, elongation and seizure pressure, if the upsetting ratio is set in a range of from 3 to 40. Contrary to the above, if the upsetting ratio exceeds 40% (in sample No. B69), cracks occur in the sample due to forging. The performance of the test in such a case has been cancelled.
Also, the effect of the heating temperature in a case where hot forging is performed is searched by comparing the samples of Nos. B68 (cold forging), B01 and B70-B75 in Tables 22 and 23 wherein that heating temperature for sintered compact is changed. From those results, it is found that, the values of tensile strength, elongation and seizure pressure are possibly improved by turning to the hot forging. This is attributable to the fact that, although in the case of cold forging hair cracks very slightly remain within the sample, followed by decrease in the elongation, carrying out hot forging of the material with the heating temperature being set to 100 degrees C. or more makes the hair cracks removed. On the other hand, when the forging temperature exceeds 400 degrees C., adhesion (die galling) of the sintered compact to the die occurs. The succeeding test in such a case has been therefore cancelled.
Also, comparing the samples of Nos.65 to 69 in Table 18, the effect of the upsetting ratio in the case where hot forging is done is searched. From those results, it is found that, in the case of hot forging, the samples have high levels of tensile strength and seizure pressure and a high value of elongation even when the upsetting ratio is extended to a wide range of 3 to 70%. However, if the upsetting ratio exceeds 70% (in sample No. B80), forging causes the occurrence of cracks on the samples. The succeeding test in such a case has been therefore cancelled.
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.
| Number | Date | Country | Kind |
|---|---|---|---|
| P2004-207586 | Jul 2004 | JP | national |
| P2005-139397 | May 2005 | JP | national |
