High-hardness au-ni-pd-pt-based precious metal alloy

Abstract

The present invention is an Au-Ni-Pd-Pt-based precious metal alloy containing 1.5-47 atom% Au, 4-57 atom% Ni, 1-43.5 atom% Pd, and 6-58.5 atom% Pt, and further containing 0.003-8 atom% B. The precious metal alloy according to the present invention achieves high hardness through a modulated structure resulting from spinodal decomposition and / or an ordered phase resulting from ordering. In addition to the high hardness, the present invention improves ductility in the recrystallized structure by adding B. Furthermore, the precious metal alloy according to the present invention can optionally include Cu, an additive metal α (In, Sn, Sb), and an additive metal β (Al, Ti, Zr, Hf).

Description

High hardness Au-Ni-Pd-Pt noble metal alloy

[0001] The present invention relates to a high-hardness precious metal alloy containing the precious metals Pt, Au, and Pd as essential constituent elements, which is designed with consideration given to workability. More specifically, the present invention relates to an Au—Ni—Pd—Pt-based precious metal alloy that has been hardened by spinodal decomposition and / or ordering, and which has improved ductility of a recrystallized structure due to solution treatment or the like, thereby improving workability.

[0002] Precious metals such as platinum (Pt) and gold (Au) have excellent chemical stability and corrosion resistance, as well as excellent electrical properties such as conductivity. Therefore, precious metals and their alloys are utilized in various fields, including electrical and electronics and medical care. Examples of applications of precious metal alloys in the electrical and electronics fields include probe pins used in semiconductor pre- and post-process testing, as well as electrical contacts (sliding contacts and open / close contacts) for motor brushes, relays, switches, and the like. Their use in the medical field has recently attracted attention, and precious metal alloys are used as components of various medical devices. Examples of such medical devices include various types of medical devices such as embolization coils, embolization clips, guidewires, stents, and catheters. These medical devices are implanted into the human body and therefore require biocompatibility and chemical stability. Furthermore, medical devices also require X-ray visibility for use in X-ray surgery and diagnosis. Precious metal alloys exhibit excellent biocompatibility and X-ray visibility.

[0003] Precious metal alloys used in the various applications described above are required to have improved mechanical properties, such as hardness and strength. For example, probe pins are required to have wear resistance because they are subjected to repeated long-term contact with mating materials. In particular, the development of probe pins with even higher hardness is required to keep up with the recent trend toward higher integration of various devices and higher performance motors. Furthermore, in the case of medical devices, devices that are implanted and moved within pulsating blood vessels, such as guide wires and embolization coils, require mechanical properties such as hardness and springiness to prevent malfunction.

[0004] Well-known methods for improving the hardness of metallic materials, including precious metal alloys, include material strengthening techniques such as work hardening (dislocation hardening), solid solution hardening, and precipitation hardening (dispersion hardening), which are used alone or in combination. Examples of methods for improving the hardness of precious metal alloys used in the above-mentioned probe pins, contact materials, medical equipment, etc. include the Pt—Ni alloy described in Patent Document 1 and the Pt—W alloy described in Patent Document 2, which achieve hardness enhancement through solid solution hardening by alloying Pt with Ni, W, etc., as well as work hardening by increasing the final work rate. Furthermore, Patent Document 3 (Ag—Pd—Cu-based alloy) and Patent Document 4 (Pt—Cr—Ni-based alloy) achieve high hardness by adjusting the work rate in addition to solid solution hardening and precipitation hardening by additive elements to obtain a precious metal alloy with high hardness.

[0005] Meanwhile, the present applicant has developed a noble metal alloy that has been hardened by spinodal decomposition and / or ordering, as a material strengthening method different from the above-mentioned well-known techniques (Patent Document 5).

[0006] Spinodal decomposition is a type of phase separation in a material structure, a phenomenon in which decomposition progresses due to a continuous increase in concentration fluctuations. The material structure produced by spinodal decomposition caused by this concentration fluctuation exhibits a very fine periodic structure of several nanometers to several tens of nanometers, called a modulated structure. In the modulated structure generated by spinodal decomposition, the concentration of solute atoms in the crystal periodically fluctuates as a function of location, and the lattice constant also changes periodically. This generates a periodic internal stress field on the slip plane, which interacts with dislocations. The strengthening mechanism by spinodal decomposition is similar to precipitation strengthening due to nucleation and growth, but differs in that it is the internal stress field generated by concentration modulation, rather than precipitates, that acts as an obstacle to dislocation movement.

[0007] Ordering is a phenomenon in which the arrangement of the constituent elements of an alloy is ordered to produce an ordered phase having a predetermined structure. The ordered phase produced by ordering contributes to increasing the hardness of the alloy due to the following factors: (i) the Burgers vector of dislocations becomes larger, (ii) antiphase boundaries can occur within the ordered phase, and (iii) the volume change associated with ordering distorts the lattice inside and outside the ordered phase, thereby suppressing dislocation motion. Ordering may occur in conjunction with the spinodal decomposition described above, or it may occur alone.

[0008] Spinodal decomposition and ordering are phenomena that can exhibit a high increase in hardness due to unique material structures such as modulated structures and ordered phases. The well-known material strengthening methods mentioned above, solid solution strengthening and precipitation strengthening, have limitations in the extent of hardness increase. Spinodal decomposition and ordering are expected to produce a higher hardening amount than these conventional strengthening methods. Furthermore, while work hardening is expected to increase hardness by increasing the degree of work, there are concerns about material embrittlement associated with hardening. Spinodal decomposition and ordering are considered to be useful means of improving hardness without inducing material embrittlement.

[0009] The above-mentioned precious metal alloy (Patent Document 5) by the present applicant is composed of an Au—Ni—Pd—Pt-based alloy, which is a quaternary or higher alloy, in order to effectively exhibit the strengthening effect due to spinodal decomposition and ordering. Binary alloys such as Pt—Au alloys and Pt—Ni alloys have been known as precious metal alloys capable of exhibiting spinodal decomposition and ordering, but their hardening amounts are not particularly high. The spinodal decomposition and ordering in these binary alloys do not surpass conventional strengthening methods such as work hardening and precipitation strengthening. The Au—Ni—Pd—Pt-based alloy by the present applicant is multi-component compared to the binary alloys and optimizes the composition ranges of each constituent element, thereby effectively exhibiting the strengthening effect due to spinodal decomposition and / or ordering. The spinodal decomposition and ordering in this Au—Ni—Pd—Pt alloy are manifested by subjecting an alloy of appropriate composition to solution treatment or the like to form a supersaturated solid solution in a recrystallized structure state, and then by subsequent aging treatment or the like, thereby achieving the desired hardness.

[0010] Japanese Patent Publication No. 2005-233967 Japanese Patent No. 6997354 Japanese Patent Publication No. 2012-242184 Japanese Patent No. 6372952 International Publication WO2023 / 063156

[0011] When precious metal alloys are used for probe pins, sliding contacts, medical equipment, etc., they must be plastically processed to the shape and dimensions required for the intended application. Regarding the workability of the Au—Ni—Pd—Pt-based alloys developed by the present applicant, these precious metal alloys have sufficient ductility and good workability even at room temperature in the cast structure or in the two-phase separated structure state that results from heat treatment at approximately 700 to 900°C. Therefore, when processing these Au—Ni—Pd—Pt-based alloys into probe pins, sliding contacts, medical equipment, etc., they are first subjected to primary processing such as rolling and wire drawing in the state with good workability (cast structure or two-phase separated structure) to produce a raw material such as a wire. The raw material after the primary processing is then solution-treated, followed by secondary processing such as skin-passing, straightening, bending, and coiling, as well as finishing, and then aging to harden it.

[0012] However, according to the investigations of the present inventors, it has been confirmed that the Au—Ni—Pd—Pt-based alloys produced by the present applicant have poor ductility after solution treatment. The decrease in ductility may affect processing, and may deteriorate the processability during the secondary processing described above.

[0013] In the Au—Ni—Pd—Pt alloy developed by the applicant, spinodal decomposition and / or ordering are performed by aging treatment to increase hardness, but in order to maximize this effect, solution treatment before aging treatment is an essential step. Therefore, when considering applying the above-mentioned processing process to the Au—Ni—Pd—Pt alloy developed by the applicant, improving ductility after solution treatment can be said to be an important issue.

[0014] The present invention has been made under the above-mentioned background, and provides a high-hardness precious metal alloy made of an Au-Ni-Pd-Pt alloy by the applicant of the present application, which exhibits hardness increased by spinodal decomposition and / or ordering, while also having improved ductility after solution treatment and the like.

[0015] The Au—Ni—Pd—Pt alloy (Patent Document 5) developed by the present applicant is an unprecedented technology in that it actively utilizes the effects of spinodal decomposition and / or ordering in precious metal alloys. Therefore, the phenomenon of ductility reduction after solution treatment as described above can be said to be unknown. In order to solve the above problem, the present inventors decided to clarify the cause of ductility reduction in Au—Ni—Pd—Pt alloys after solution treatment.

[0016] As a result, the present inventors discovered that Au-Ni-Pd-Pt alloys, which can undergo spinodal decomposition and ordering, can have grain boundary segregation of Au and / or Pd in ​​the recrystallized structure, which is a material structure formed by solution treatment. The present inventors then considered that this grain boundary segregation causes grain boundary embrittlement in the recrystallized structure, which is the cause of the decrease in ductility. Note that such a recrystallized structure is a material structure that is observed in all or part of the Au-Ni-Pd-Pt alloy after solution treatment or the like.

[0017] If the cause of the decrease in ductility after solution treatment is assumed to be grain boundary embrittlement due to the segregation of Au and / or Pd in ​​the recrystallized structure, the prerequisite for addressing this issue is to improve ductility within a range that does not inhibit the hardening effects of spinodal decomposition and ordering. In other words, even if the segregation of Au and Pd is to be suppressed, it is not desirable to consider changing the composition of Au, Pd, and other constituent elements to a range in which neither spinodal decomposition nor ordering occurs. Therefore, the present inventors attempted to search for an element in an Au-Ni-Pd-Pt-based alloy that segregates to grain boundaries and strengthens the grain boundaries without inhibiting spinodal decomposition and ordering. As a result, they found that B (boron) satisfies the above conditions, leading to the invention.

[0018] That is, the present invention is an Au-Ni-Pd-Pt-based precious metal alloy containing 1.5 atomic % or more and 47 atomic % or less of Au, 4 atomic % or more and 57 atomic % or less of Ni, 1 atomic % or more and 43.5 atomic % or less of Pd, and 6 atomic % or more and 58.5 atomic % or less of Pt, and the Au-Ni-Pd-Pt-based precious metal alloy contains 0.003 atomic % or more and 8 atomic % or less of B.

[0019] The precious metal alloy according to the present invention contains Au, Ni, Pd, Pt, and B as essential constituent elements, and is allowed to contain some additional elements, and can contain more than 0.1 atomic % and not more than 27.5 atomic % of Cu.

[0020] Furthermore, the Au—Ni—Pd—Pt-based precious metal alloy according to the present invention may contain, as optional additive elements, 0.15 atomic % to 5 atomic % of a metal element α and / or 0.05 atomic % to 5 atomic % of a metal element β, where the metal element α is at least one of In, Sn, and Sb, and the metal element β is at least one of Al, Ti, Zr, and Hf.

[0021] As described above, the Au-Ni-Pd-Pt based noble metal alloy according to the present invention contains a modulated structure and / or an ordered phase due to spinodal decomposition.

[0022] The Au—Ni—Pd—Pt-based noble metal alloy according to the present invention is hardened by spinodal decomposition and / or ordering, and the hardened Au—Ni—Pd—Pt-based noble metal alloy has a Vickers hardness of 500 Hv or more.

[0023] As described above, the Au—Ni—Pd—Pt-based precious metal alloy according to the present invention has improved ductility in the recrystallized structure after solution treatment. As with the Au—Ni—Pd—Pt-based precious metal alloy developed by the present applicant described above, the present invention can achieve high hardness by a modulated structure due to spinodal decomposition and / or an ordered phase due to ordering, which are developed by appropriate heat treatment, such as aging treatment. According to the present invention, even in the recrystallized structure state after solution treatment, workability into shapes suitable for various applications can be ensured, and suitable hardness can be achieved by subsequent aging treatment.

[0024] 1 shows the results of XRD of a solution-treated material and an aging-treated material of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2. 2 shows a STEM-EDS mapping image of a sample of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 that was subjected to aging treatment after solution treatment. 3 shows an electron diffraction pattern of a sample of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 that was subjected to aging treatment after solution treatment. Photographs showing the results of STEM-EDS analysis showing the element distribution near high-angle grain boundaries of the precious metal alloy (Au7.5-Ni41.25-Pd10-Pt41.25) of Comparative Example 2. A diagram showing Auger spectra of the grain boundary fracture surface and intragranular fracture surface after impact fracture of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2. A diagram showing the results of room temperature tensile tests of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 and the precious metal alloy (Au12.5-Ni37.5-Pd12.5-Pt37.5) of Comparative Example 4-1.

[0025] The following describes embodiments of the present invention. As described above, the precious metal alloy according to the present invention is an Au—Ni—Pd—Pt alloy in which B, an additive element, is added to an Au—Ni—Pd—Pt alloy. The configuration of the precious metal alloy according to the present invention and a method for producing the precious metal alloy according to the present invention will be described below.

[0026] (A) Structure of the Precious Metal Alloy According to the Present Invention (A-1) Constituent Elements and Composition Ranges of the Precious Metal Alloy According to the Present Invention The precious metal alloy according to the present invention contains Au, Ni, Pd, and Pt as essential constituent elements, and also contains B as an essential additive element. In the following explanation, the essential constituent elements and their composition ranges will be clarified. Then, optional additive elements (Cu, metal element α, and metal element β) that are permitted to be added in addition to B in the precious metal alloy according to the present invention and their addition concentrations will also be explained.

[0027] (1) Essential Constituent Elements (Au, Ni, Pd, Pt) The precious metal alloy according to the present invention contains the above essential elements, and thereby exhibits at least one of spinodal decomposition and ordering, and has a hardness equal to or greater than that of precious metal alloys formed by conventional general strengthening mechanisms (solid solution strengthening, precipitation strengthening, work hardening).

[0028] As explained in Patent Document 5, which discloses an Au—Ni—Pd—Pt alloy by the applicant of the present application, periodic concentration fluctuations occur in the modulated structure formed by spinodal decomposition, which contributes to high hardness by forming an internal stress field around it. The resistance force (critical shear stress) to dislocation motion in this periodic internal stress field is expressed by the following formula, and the lattice strain (ε), elastic modulus (Y), and concentration modulation amplitude (A) are thought to be the governing factors (for detailed references, see, for example, "Introduction to Dislocation Theory" by Masaharu Kato (published in August 1999, published by Shokabo)).

[0029]

[0030] In light of the above formula, it is preferable that the constituent elements of a precious metal alloy that effectively exhibits material strengthening through spinodal decomposition satisfy the following three requirements: (1) the constituent elements include a metal with a high elastic modulus, (2) the mixing enthalpy between the constituent elements is high and there is a strong tendency for phase separation in the low temperature range, and (3) the difference in lattice constant between the constituent elements is large and the lattice strain (ε) is also large.

[0031] Regarding the occurrence of ordering, Pt and Ni, and Au and Pd are combinations of metals that can contribute to the generation of an ordered phase through ordering.

[0032] The essential elements of the precious metal alloy of the present invention, Au, Ni, Pd, and Pt, are a combination of metals that can satisfy these requirements, and by setting the composition ranges described below, the strengthening ability by spinodal decomposition works effectively. The specific functions and composition ranges of these essential elements are explained below.

[0033] Au Au is an essential element for causing spinodal decomposition in the alloy system of the present invention. Spinodal decomposition does not occur if the Au concentration is too low or too high, and there is a concentration range of Au necessary for its occurrence. If the Au concentration is outside the optimal range, normal nucleation and growth tend to occur, making it impossible to obtain a suitable increase in hardness. In addition, Au is a metal that can form an ordered phase with Pd, and also has the effect of contributing to the increase in hardness due to ordering.

[0034] The Au concentration in the Au—Ni—Pd—Pt alloy according to the present invention is 1.5 atomic % or more and 47 atomic % or less, preferably 4 atomic % or more and 30 atomic % or less, more preferably 6 atomic % or more and 24 atomic % or less, and even more preferably 7 atomic % or more and 20 atomic % or less.

[0035] Ni Ni acts as a strengthening factor when spinodal decomposition occurs in precious metal alloys. Ni has a higher elastic modulus than Au, Pt, and Pd. In addition, the lattice constant of Ni is smaller than those of Au, Pt, and Pd, so it has a large effect of increasing the lattice strain ε. Therefore, from the above equation 1, Ni has the effect of increasing the strengthening ability due to spinodal decomposition. Furthermore, Ni is a metal that can form an ordered phase with Pt, and also has the effect of contributing to an increase in hardness due to ordering.

[0036] In addition, Ni is an element of the same group as Pt and Pd and has a similar electronic structure, so it can be used to form alloys without impairing the corrosion resistance and oxidation resistance of precious metals to the greatest extent possible, which also has the secondary effect of reducing the price of the entire precious metal alloy.

[0037] The Ni concentration of the Au—Ni—Pd—Pt alloy according to the present invention is 4 atomic % or more and 57 atomic % or less, preferably 10 atomic % or more and 52.5 atomic % or less, more preferably 12.5 atomic % or more and 50 atomic % or less, and even more preferably 15 atomic % or more and 47.5 atomic % or less.

[0038] Pd Pd has the effect of expanding the solid solubility limits of each element constituting the precious metal alloy, widening the concentration range in which spinodal decomposition of the precious metal alloy occurs, and promoting spinodal decomposition. As a result, Pd has the effect of improving the hardening amount due to spinodal decomposition. Furthermore, Au is a metal that can form an ordered phase with Pd, and also contributes to increasing hardness through ordering. However, if excessive Pd is added, the spinodal decomposition temperature will excessively decrease, which will tend to inhibit spinodal decomposition. Furthermore, excessive addition of Pd will tend to suppress ordering, resulting in a decrease in the total hardening amount of the alloy system. Therefore, in order to optimize the hardening amount of the precious metal alloy, there is an optimal concentration range for Pd as described above.

[0039] The Pd concentration of the Au—Ni—Pd—Pt alloy according to the present invention is 1 atomic % or more and 43.5 atomic % or less, preferably 3 atomic % or more and 37.5 atomic % or less, more preferably 4 atomic % or more and 32.5 atomic % or less, and even more preferably 5 atomic % or more and 27.5 atomic % or less.

[0040] Pt Pt is also an essential element for spinodal decomposition in the alloy system of the present invention. Spinodal decomposition does not occur if the Pt concentration is too low or too high, and there is a Pt concentration range required for its occurrence. Pt also forms an ordered phase with Ni, which can contribute to an increase in hardness. Furthermore, since Pt has a relatively high elastic modulus, it can be expected to be an element that increases the hardness of the alloy when spinodal decomposition occurs.

[0041] The Pt concentration in the Au—Ni—Pd—Pt alloy according to the present invention is 6 atomic % or more and 58.5 atomic % or less, preferably 12.5 atomic % or more and 52.5 atomic % or less, more preferably 15 atomic % or more and 50 atomic % or less, and even more preferably 17.5 atomic % or more and 47.5 atomic % or less.

[0042] (2) Essential Additional Element (B) B, an essential additional element in the Au—Ni—Pd—Pt-based alloy according to the present invention, is an element that exerts grain boundary strengthening of the precious metal alloy by segregating at the grain boundaries of the recrystallized structure formed by solution treatment, etc. This suppresses the reduction in ductility or embrittlement that may occur due to the grain boundary segregation of Au and Pd at the site where the recrystallized structure has formed in the precious metal alloy after solution treatment, and imparts suitable workability to the precious metal alloy containing the recrystallized structure.

[0043] The B concentration in the precious metal alloy according to the present invention is 0.003 atomic % or more and 8 atomic % or less. If it is less than 0.003 atomic %, the amount of segregation to the grain boundaries of the recrystallized structure is small, and the above-mentioned grain boundary strengthening effect cannot be fully obtained. Furthermore, if it exceeds 8 atomic %, ductility tends to decrease not only in the recrystallized structure state but also in the cast structure and dual-phase structure state, which is undesirable. The B concentration is preferably 0.0075 atomic % or more. Furthermore, the B concentration can be appropriately set to 0.01 atomic % or more and 4 atomic % or less, more than 0.1 atomic % and 3 atomic % or less, or 0.2 atomic % or more and 2.5 atomic % or less, taking into consideration the overall alloy composition of the Au—Ni—Pd—Pt-based alloy.

[0044] (3) Optionally Added Elements (Cu, Metal Element α, Metal Element β) The Au—Ni—Pd—Pt alloy according to the present invention has a base alloy (base alloy) made of a precious metal alloy containing the above-mentioned essential constituent elements, and may contain optional additional elements. Examples of the optional additional elements include Cu, metal element α, and metal element β.

[0045] Cu Cu has the effect of expanding the solid solubility limit between the constituent elements of a precious metal alloy, thereby broadening the concentration range for the occurrence of spinodal decomposition. In this respect, Cu has a similar effect to Pd and can therefore be used as an additive element. However, even if some of the effects of Pd are similar, Cu is an optional additive element and is not an essential constituent element like Pd. This is because Cu has the effect of broadening the concentration range for the occurrence of spinodal decomposition, but does not have the effect of improving the hardening amount due to spinodal decomposition. In addition to the above effects, Cu as an optional additive element has the effect of improving machinability, and further, since it is a cheaper metal than precious metals, it has the secondary effect of reducing the price of the entire precious metal alloy.

[0046] When adding Cu, an optional element, to obtain the above-mentioned effects, the concentration is preferably more than 0.1 atomic % and less than 27.5 atomic %. Excessive Cu addition can lower the spinodal decomposition temperature, suppress ordering, and potentially reduce the total hardening amount of the alloy system. Excessive Cu addition can also cause a decrease in the corrosion resistance of the alloy. When Cu is added, the Cu concentration is more preferably 2.5 atomic % to 22.5 atomic % and particularly preferably 5 atomic % to 18.5 atomic %.

[0047] Metal Element α, Metal Element β The Au—Ni—Pd—Pt alloy according to the present invention may contain, in addition to the essential constituent elements and the above-mentioned Cu, metal elements α and β as optional additive metal elements. Metal elements α and β have the effect of increasing the hardness of the precious metal alloy after aging treatment. While the reason why the addition of metal elements α and β increases the hardness of the precious metal alloy is unclear, the inventors believe that it may be related to factors such as increased lattice strain resulting in increased strength of the alloy matrix, refinement of the alloy crystal grains after aging treatment, and the formation of precipitates containing the additive elements. The addition of metal elements α and β does not inhibit the formation of a material structure containing a modulated structure and / or an ordered phase. Either or both of metal elements α and β may be included.

[0048] The metal element α is at least one of In, Sn, and Sb. According to the studies of the present inventors, these metal elements have been found to increase the hardness of the precious metal alloy of the present invention. It is preferable to add one or more of In, Sn, and Sb as the metal element α. The composition range of at least one of In, Sn, and Sb, which is the metal element α, is 0.15 atomic % or more and 5 atomic % or less. An addition amount of less than 0.15 atomic % is unlikely to contribute to increasing the hardness of the precious metal alloy, and an addition amount exceeding 5 atomic % significantly reduces workability. The composition range of the metal element α is preferably 0.3 atomic % or more and 3 atomic % or less, and more preferably 0.5 atomic % or more and 1.5 atomic % or less.

[0049] The metal element β is at least one of Ti, Zr, Hf, and Al. According to the inventors' studies, these metal elements have been found to refine the crystal grains and increase the hardness of the precious metal alloy of the present invention. The metal element β is preferably one or more of Ti, Zr, Hf, and Al. The composition range of at least one of Ti, Zr, Hf, and Al, which is the metal element β, is 0.05 atomic % or more and 5 atomic % or less. An addition amount of less than 0.05 atomic % does not contribute to refinement of the crystal grains or increase in hardness of the precious metal alloy, and an addition amount exceeding 5 atomic % significantly reduces workability. The composition range of the metal element β is preferably more than 0.1 atomic % and 3 atomic % or less, and more preferably 0.15 atomic % or more and 1.5 atomic % or less.

[0050] The precious metal alloy according to the present invention is a precious metal alloy containing Au, Ni, Pd, Pt, and B as essential elements within the above composition ranges, and is preferably a precious metal alloy consisting of Au, Ni, Pd, Pt, and B within the above composition ranges. However, the precious metal alloy according to the present invention may contain any unavoidable impurities. These unavoidable impurities are unavoidable components that are present due to impurities in the raw materials or the manufacturing process. Specific examples of unavoidable impurities include Rh, Ir, Fe, Y, Zn, Co, Si, Cu, Th, H, and rare earth elements. These unavoidable impurities are introduced from raw materials and equipment used during melt casting. The content of these unavoidable impurities is preferably within a range that does not impair the properties of the precious metal alloy according to the present invention, and is preferably less than 0.05 atomic % per element, preferably 0.5 atomic % or less in total, and particularly preferably less than 0.1 atomic % in total. In addition, when the above-mentioned inevitable impurities are contained in a precious metal alloy, it is difficult to clearly distinguish whether they are components that are inevitably contained or components that are intentionally added. In the present invention, as long as the components do not significantly change the properties of the precious metal alloy, they are considered to be inevitable impurities regardless of the intention of their inclusion. The above definition of inevitable impurities also applies to the case where the above-mentioned additive elements Cu, metal element α, and metal element β are contained.

[0051] The spinodal decomposition and ordering of the precious metal alloy according to the present invention are manifested by a solution treatment in which a solid solution alloy having the above-described composition range is held at high temperature and then rapidly cooled, followed by a subsequent aging treatment. The precious metal alloy according to the present invention has a wide range of complete solid solution at high temperatures, while having a solubility gap at low temperatures. Therefore, it is believed that a supersaturated solid solution is formed when a precious metal alloy having the above-described composition range is solution-treated at high temperatures and then rapidly cooled, and spinodal decomposition and ordering can occur by subsequent aging treatment. Regarding the thermodynamic behavior (transformation point, phase equilibrium, solid solubility limit, melting point, etc.) of the precious metal alloy according to the present invention, the CALPHAD method (Calculation of Phase Diagrams) can also be effectively used. For calculations using the CALPHAD method, it is preferable to use commercially available thermodynamic calculation software (e.g., Thermo-Calc and a precious metal alloy database (e.g., TCNOBL1)).

[0052] (A-2) Material structure of the precious metal alloy according to the present invention The precious metal alloy according to the present invention has increased hardness due to spinodal decomposition and / or ordering. Therefore, the material structure of the precious metal alloy according to the present invention may include a finely modulated structure due to spinodal decomposition and / or an ordered phase due to ordering.

[0053] In the precious metal alloy according to the present invention, the modulated structure caused by spinodal decomposition is a material structure whose composition fluctuates with a modulation period of the order of several nanometers to several tens of nanometers. Specifically, the modulated structure in the present invention is composed of two regions: a region where the Au and Pd concentrations are relatively high (a region where the Pt and Ni concentrations are relatively low) and a region where the Au and Pd concentrations are relatively low (a region where the Pt and Ni concentrations are relatively high).

[0054] Furthermore, the composition and structure of the ordered phase in the present invention are not necessarily completely clear. However, it is believed that a phase having the same or similar crystal structure as the ordered phase that can be formed in Pt—Ni based alloys or Au—Pd based alloys, which are known as noble metal alloys in which ordering can occur, is also formed in the noble metal alloy of the present invention. Therefore, the ordered phase in the present invention is considered to be L1 ０ Type structure or L1 ２It is presumed to be a type structure or a similar crystalline structure.

[0055] The material structure of the precious metal alloy of the present invention can be confirmed by X-ray diffraction (XRD), transmission electron microscope (TEM), TEM electron diffraction pattern, scanning transmission electron microscope (STEM), TEM / STEM-EDS mapping image, TEM / STEM-EELS mapping image, three-dimensional atom probe, etc.

[0056] The modulated structure of the precious metal alloy of the present invention can be confirmed by one or both of diffraction method and elemental analysis (elemental mapping).

[0057] Diffraction methods include confirming X-ray diffraction patterns by X-ray diffraction (XRD) or electron beam diffraction patterns by TEM. When confirmed by diffraction, if a modulated structure is present, so-called sideband peaks (satellite peaks) or broad peaks called streaks are observed on at least one side (preferably both sides) of the main peak or regular reflection peak. The presence or absence of these sideband peaks can determine whether a modulated structure due to spinodal decomposition is present. For example, in a diffraction pattern obtained by X-ray diffraction, since the crystal structure of the matrix of the precious metal alloy according to the present invention is a face-centered cubic (fcc) structure, Miller index planes {111}, {200}, {220}, {311}, etc. appear as main peaks. In the case of a spinodal decomposition structure, sideband peaks appear on both sides or one side of at least one of the main peaks. Note that if a sideband peak appears only on one side of the main peak, it is thought that it is difficult to separate it from the main peak.

[0058] On the other hand, confirmation methods using element mapping include methods that visually confirm the density of atomic concentrations using STEM-EDS / EELS or a three-dimensional atom probe.

[0059] The presence or absence of ordered reflection peaks is generally confirmed in an X-ray diffraction pattern or a TEM electron diffraction pattern. For example, in observing ordered reflection peaks in an X-ray diffraction pattern, in the case of θ-2θ measurement using CuKα radiation as the X-ray source, ordered reflection peaks appear around 2θ=22.5 to 27.5°, around 30 to 35°, etc. In addition, in the case of an electron diffraction pattern, diffraction spots such as 100, 110, and 120 appear, which do not appear in an fcc structure.

[0060] (A-3) Hardness of the Precious Metal Alloy According to the Present Invention The precious metal alloy according to the present invention has the composition range described above, and is hardened by spinodal decomposition and / or ordering. The precious metal alloy according to the present invention can stably exhibit a Vickers hardness of 500 Hv or more. The Vickers hardness of the precious metal alloy according to the present invention is preferably 540 Hv or more, more preferably 590 Hv or more, and even more preferably 640 Hv or more. The precious metal alloy according to the present invention can be made into a high-hardness precious metal alloy having the above-mentioned Vickers hardness by heat treatment alone, without utilizing work hardening at all, i.e., without causing material embrittlement due to dislocation strain.

[0061] The upper limit of the hardness of the precious metal alloy according to the present invention should not be particularly limited, but it is preferable that the upper limit is 850 Hv or less. If the hardness exceeds 850 Hv, there is a risk of breakage or chipping during use. The Vickers hardness described above is a value at room temperature. The Vickers hardness can be measured using a known Vickers hardness tester. The measurement load is preferably 0.025 kgf or more and 0.5 kgf or less, and more preferably 0.075 kgf.

[0062] The shape and form of the precious metal alloy according to the present invention are not particularly limited. For the above-mentioned medical instruments, probe pins, etc., the present invention can generally be used as a bulk (lump) alloy that has been appropriately processed. The present invention has improved ductility so that it can be processed into shapes and dimensions suitable for these applications. The present invention can also be formed in the form of a layer or film on an appropriate base material or substrate.

[0063] (B) Manufacturing Method of Precious Metal Alloy According to the Present Invention Next, a method for manufacturing a precious metal alloy according to the present invention will be described. As described above, the precious metal alloy according to the present invention can be provided in various shapes and forms. In the following description, a method for manufacturing a bulk precious metal alloy, which is a frequently applied form, will be described in detail, and a method for manufacturing a layered or film-shaped precious metal alloy will also be mentioned.

[0064] In the present invention, high hardness is achieved by selecting the constituent elements of the precious metal alloy (essential constituent elements (Au, Ni, Pd, Pt, B) and optionally added metal elements (Cu, metal elements α, β)) and optimizing their composition ranges, and then causing spinodal decomposition and / or ordering. The precious metal alloy of the present invention is produced by producing and preparing an alloy block (ingot) having the above-mentioned composition range, and then carrying out an appropriate heat treatment step. The appropriate heat treatment step is a heat treatment step that combines solution treatment and aging treatment, and by undergoing these steps, spinodal decomposition and / or ordering proceed, resulting in high hardness and the precious metal alloy of the present invention.

[0065] The method for producing a precious metal alloy according to the present invention will be described below with reference to the respective heat treatment steps of solution treatment and aging treatment. In the present invention, the temperatures, such as the heating temperature, in the various heat treatments described below refer to the temperature of the precious metal alloy being treated, unless otherwise specified.

[0066] (B-1) Preparation Step (Preparation of Precious Metal Alloy) First, an unheated precious metal alloy ingot is prepared as a precursor for the precious metal alloy of the present invention. The precursor precious metal alloy ingot can be produced by a conventional melt casting method. The above-mentioned metal raw materials, Au, Ni, Pd, Pt, B, and optional elements Cu, and metallic elements α and β, are adjusted to the above composition by appropriate weighing or the like, and then melted and cast to produce an alloy ingot. In this case, a precious metal alloy not containing Cu or metallic elements α and β (Au—Ni—Pd—Pt alloy), or a binary alloy such as an Au—Pd alloy or a Pt—Ni alloy may be appropriately combined and melted as a master alloy. Melt casting of the precious metal alloy can be performed by known methods such as arc melting, high-frequency melting, vacuum melting, continuous casting, and liquid quenching.

[0067] The precious metal alloy may be prepared by a method other than melt casting, such as powder metallurgy. In powder metallurgy, a precious metal alloy powder prepared according to the above composition (e.g., a precious metal alloy powder produced by atomization) is sintered to obtain an alloy ingot to be heat-treated. Alternatively, a near-net-shape ingot may be produced using a known metal powder injection molding method or additive manufacturing method using the precious metal alloy powder prepared according to the above composition. Furthermore, a precious metal alloy layer of the above composition may be formed on any base material by a known alloy formation method, such as plating, sputtering, or thermal spraying.

[0068] (B-2) Solution Treatment: The precious metal alloy prepared as described above is subjected to solution treatment to form a supersaturated solid solution. Solution treatment is a process in which a precious metal alloy is heated to a high temperature to form a structure with a high solid solution concentration, followed by rapid cooling to form a supersaturated solid solution. The heating temperature for solution treatment is preferably at or above (Tm - 500°C) but below Tm, where Tm is the melting point (solidus) of the precious metal alloy (°C). Temperatures below (Tm - 500°C) result in low solid solubility of each element and are insufficient to form a supersaturated solid solution. Temperatures above Tm are undesirable because melting of the material begins near the grain boundaries. The holding time during heating is preferably in the range of 0.0001 to 168 hours. Holding times of less than 0.0001 hours result in insufficient formation of a supersaturated solid solution, and heating for 168 hours or more does not significantly affect the formation of a supersaturated solid solution, making this undesirable from the standpoint of productivity. In the present invention, the melting point means the solidus temperature.

[0069] Furthermore, cooling from the solution treatment temperature requires rapid cooling, i.e., rapid cooling, so as not to cause excessive grain boundary reactions in the high-temperature range. If grain boundary reactions occur in the high-temperature range, the grain boundaries are strengthened, improving ductility, but the hardness after aging may be insufficient. Specifically, the cooling rate is preferably 10°C / s or higher, more preferably 50°C / s or higher, and even more preferably 150°C / s or higher. On the other hand, a slow cooling rate is preferable from the viewpoint of preventing quench cracking, dimensional change, deformation, etc. Therefore, from the viewpoint of improving the hardness of the precious metal alloy, the cooling rate referred to as rapid cooling is not necessary in the low-temperature range where grain boundary reactions do not occur and spinodal decomposition and / or ordering do not proceed excessively. For example, rapid cooling is not required for the cooling rate in the temperature range below 250°C. Therefore, to suppress or reduce the occurrence of quench cracking, etc., rapid cooling to 200°C may be performed, followed by air cooling in the temperature range below 200°C. The end point of cooling in the solution treatment is preferably room temperature.

[0070] (B-3) Aging Treatment Spinodal decomposition and ordering in the precious metal alloy according to the present invention proceed by aging the supersaturated solid solution formed by solution treatment at a temperature lower than the spinodal decomposition temperature and the order-disorder transformation temperature.

[0071] As conditions for the aging treatment of a supersaturated solid solution precious metal alloy, the heating temperature is preferably 250°C or higher and 800°C or lower. If the temperature is lower than 250°C, spinodal decomposition or ordering is unlikely to proceed. If the temperature exceeds 800°C, significant softening of the material due to grain boundary reaction occurs. The heating temperature is more preferably 350°C or higher and 650°C or lower. The heating time for the aging treatment is preferably 0.001 hours or higher and 168 hours or lower. If the heating time is less than 0.001 hours, the transformation is insufficient, resulting in variations in hardness, while treatment for 168 hours or longer results in poor productivity and high manufacturing costs. There are no particular restrictions on the cooling method after the aging treatment. The high-hardness precious metal alloy according to the present invention can be obtained by undergoing this aging treatment.

[0072] (B-4) Other Heat Treatment Steps In the production of the precious metal alloy according to the present invention, the above-mentioned solution treatment and aging treatment are essential steps, but other heat treatment steps may also be included. Examples of such other heat treatment steps include homogenization treatment, two-phase treatment, intermediate annealing, etc. However, these other heat treatment steps do not affect the progress of spinodal decomposition or ordering. Therefore, these heat treatment steps are optional steps.

[0073] Homogenization is a heat treatment performed on precious metal alloys prepared by melt casting to form a metal structure with a uniform distribution of element concentrations in the precious metal alloy. The homogenization is a heat treatment in which the precious metal alloy is heated at a high temperature below its melting point for a long period of time (preferably 0.1 to 72 hours).

[0074] The dual-phase treatment is a heat treatment for forming a two-phase structure, which has the best workability among the alloys of the present invention, and is carried out to facilitate warm working and cold working. The heating temperature for the dual-phase treatment is 700°C or higher and 900°C or lower, more preferably 750°C or higher and 850°C or lower. The heating time is 0.1 hours or higher and 10 hours or lower, more preferably 0.2 hours or higher and 2 hours or lower.

[0075] Intermediate annealing is a heat treatment performed during processing of precious metal alloy ingots, etc., in which strain accumulates during warm working or cold working, as described below. Intermediate annealing is a treatment for reducing the material strength and restoring workability. The heating temperature for intermediate annealing is 700°C or higher and 900°C or lower, more preferably 750°C or higher and 850°C or lower. The heating time is 0.1 hours or higher and 10 hours or lower, more preferably 0.2 hours or higher and 2 hours or lower.

[0076] (B-5) Processing Process of the Precious Metal Alloy According to the Present Invention The precious metal alloy according to the present invention can be processed into various shapes depending on the application by undergoing at least one processing step before aging treatment. Examples of this processing step include hot working, warm working, cold working, skin-pass, straightening, coiling, bending, etc. Hot working can destroy the solidification structure in the prepared precious metal alloy ingot and eliminate defects such as voids. Warm working and cold working not only change the overall shape of the alloy, but also have the significance of controlling the shape of the crystal grains. When warm working or cold working is performed multiple times, the intermediate annealing described above may be performed between processing passes.

[0077] The precious metal alloy according to the present invention, which has improved ductility in the recrystallized structure, exhibits its effects particularly when the solution-treated precious metal alloy is subjected to warm working, cold working, skin-pass processing, straightening, coiling, and bending. It is expected that after the solution treatment (before aging), the precious metal alloy will be worked into a final shape or a shape close to the final shape for its intended use. The precious metal alloy according to the present invention has improved ductility due to grain boundary segregation and grain boundary strengthening of the added element B, and can be worked without defects even in the recrystallized structure state after the solution treatment. It should be noted that the precious metal alloy according to the present invention can be worked after aging (hardening). Even a precious metal alloy with high hardness after aging can still be worked. Furthermore, after aging, final adjustment processes such as grinding / polishing, cutting, electric discharge machining, pressing, bending, and straightening are also possible. Furthermore, because the spinodal decomposition and ordering exhibited in the precious metal alloy according to the present invention are reversible, the solution treatment and aging treatment may be repeated. Working may be carried out between a combination of multiple solution treatments and aging treatments.

[0078] (B-6) Other Manufacturing Methods of the Precious Metal Alloy According to the Present Invention In the manufacturing methods described above, the precious metal alloy according to the present invention is obtained by subjecting an alloy ingot that has undergone a melting and casting process to a solution treatment and aging treatment. In the present invention, the solution treatment is the main heat treatment, but in some cases, a precious metal alloy having a structure equivalent to that of a supersaturated solid solution with a recrystallized structure can be prepared without performing the solution treatment step. In such cases, the precious metal alloy according to the present invention can be manufactured by subjecting the precious metal alloy to an aging treatment. Furthermore, the addition of B according to the present invention can contribute to improving the ductility of the precious metal alloy in such a recrystallized structure or supersaturated solid solution structure before the aging treatment.

[0079] Examples of processes for obtaining a solid solution state of a precious metal without solution treatment include a liquid quenching method applied to molten precious metal alloys, and a rapid heating and quenching process or rapid solidification process using a laser or the like applied to bulk metals. Furthermore, various film-forming processes, such as welding, sputtering, plating, and thermal spraying, can also produce precious metal alloys that approximate a supersaturated solid solution state. Precious metal alloys produced by these processes can be obtained as the precious metal alloy of the present invention by aging treatment without solution treatment. These processes are useful for producing near-net-shape precious metal alloys and layered or film-like precious metal alloys for high-hardness coatings, etc. Even when such solution treatment is not performed, the preferred conditions for aging treatment are the same as those described above.

[0080] Specific examples of the present invention are described below. In these examples, multiple Au—Ni—Pd—Pt-based precious metal alloys were manufactured by varying the composition of each element: Au, Ni, Pd, Pt, B, Cu, metal element α, and metal element β. The manufactured precious metal alloys were then evaluated for ductility in a state after solution treatment (recrystallized structure state), and their hardness after aging treatment was measured. The manufacturing process and evaluation method for the precious metal alloys of each example are as follows.

[0081] [Production of precious metal alloy] High-purity raw materials of Au, Ni, Pd, Pt, B, Cu, and metal elements α and β were weighed and mixed to obtain a predetermined composition, and then placed in an alumina crucible. After evacuation, the crucible was subjected to high-frequency melting in a reduced-pressure argon atmosphere and cast into a copper mold (diameter 11 mm, length 70 mm). The resulting ingot was then subjected to homogenization treatment. The homogenization treatment was carried out by setting the heating temperature between 800°C and 1225°C, taking into consideration the composition and melting point of the precious metal alloy, and heating the mixture in a vacuum atmosphere (5 × 10 ―２ The heating time for the homogenization treatment was 16 hours in all cases.

[0082] The homogenized ingots (11 mm in diameter) were subjected to a two-phase treatment (800°C for 60 minutes in a vacuum atmosphere), followed by cold working to a 6 mm diameter by swaging and then cold working to a 2.5 mm square by groove rolling. In these processing steps, the processing rate per pass was set to 10-15%, and intermediate annealing (800°C for 60 minutes in a vacuum atmosphere) was performed within a total processing rate range of 30-50%. Furthermore, wire drawing was performed in a wire drawing machine at a processing rate of 10-20% per pass to a 0.6 mm diameter. In this wire drawing, intermediate annealing was performed within a total processing rate range of 30-60%. Through these processing steps, the precious metal alloys of each Example and Comparative Example were processed into wire rods with a diameter of 0.6 mm. These wire rods were used as samples for evaluating the ductility of the recrystallized state and measuring hardness.

[0083] [Solution Treatment] The wire samples of each example were subjected to solution treatment in a vacuum atmosphere at a heating temperature of 925°C to 1225°C for a heating time of 1 minute to 5 minutes. After heating, the samples were placed in a water-cooled bath (20±5°C) within 3 seconds to cool to room temperature.

[0084] [Aging Treatment] After the solution treatment, each wire rod sample was subjected to a bending test to evaluate ductility, followed by aging treatment. The bending test method will be described later. The heating temperature during the aging treatment was set between 450°C and 600°C, and the samples were heated and held at each aging treatment temperature for 1 hour, followed by water cooling to room temperature. After the aging treatment, the sample pieces were embedded in resin and roughly polished (#500, #800, #1200) and mirror-polished with a 1 μm and 1 / 4 μm diamond suspension to remove residual stress in the surface due to oxidation layers and thermal strain and to prepare samples for hardness measurement.

[0085] In this embodiment, Au—Ni—Pd—Pt-based precious metal alloys were manufactured as examples, and Au—Ni—Pd—Pt alloys with compositions similar to those of the precious metal alloys of the examples but without the addition of B were manufactured as comparative examples. Furthermore, Au—Ni—Pd—Pt-based precious metal alloys with compositions similar to those of the examples but with C (carbon) added instead of B were also manufactured as comparative examples. Furthermore, Au—Ni—Pd—Pt-based precious metal alloys with Ag and Zr added were manufactured as comparative examples related to the grain boundary strengthening effect due to grain boundary segregation.

[0086] [Study of Material Structure by XRD Analysis] Furthermore, XRD analysis was performed on each precious metal alloy after aging treatment to confirm the occurrence of spinodal decomposition and the formation of an ordered phase. The XRD analysis conditions for each investigation were as follows. XRD analysis was performed on the precious metal alloy after solution treatment (solution-treated material) and the precious metal alloy after aging treatment after solution treatment (aging-treated material). By comparing the results, it was possible to confirm the occurrence of spinodal decomposition and ordering due to aging treatment. Note that the samples for XRD analysis were prepared by casting precious metal alloys with a diameter of 30 mm and a length of 100 mm, subjecting them to the same homogenization treatment and two-phase treatment (800°C, 60 minutes) as described above, cold swaging to a diameter of 22 mm, and then subjecting them to the same solution treatment (holding time: 10 minutes) and aging treatment as described above, in order to improve the S / N ratio of ordered reflection peaks with low diffraction intensity. The samples after the solution treatment (solution treated material) and the samples after the aging treatment (aging treated material) were subjected to XRD analysis. The analysis samples were embedded in resin and polished in the same manner as described above.

[0087] [Common conditions] Sample size: φ22 mm x 2 mmt XRD device: Rigaku SmartLab Target: Cu anode Optical system / detector: Concentrated optical system / semiconductor detector (HyPix-3000) Voltage / current: 40 kV / 30 mA Length limiting slit: 10 mm

[0088] (a) Confirmation of spinodal decomposition (sideband peak) 2θ scanning range: 20° to 130° 2θ step size (°): 0.0012 2θ scanning rate (° / min): 6

[0089] (b) Confirmation of ordered phase (ordered peak) 2θ scanning range: 20° to 38° 2θ step size (°): 0.0132 2θ scanning rate (° / min): 1.3

[0090] In examining spinodal decomposition by XRD, the XRD diffraction profile obtained under the above conditions was judged based on whether or not one or more side band peaks appeared on either or both sides (approximately ±0.5 to 3° in 2θ angle) of the main peak for the Miller indices {111}, {200}, {220}, and {311} planes. When one or more side band peaks appeared, it was deemed that spinodal decomposition had occurred, and when no side band peaks appeared, it was deemed that spinodal decomposition had not occurred, and evaluation was performed.

[0091] In addition, the examination of the ordered phase by XRD was judged based on the presence or absence of an ordered reflection peak occurring in the vicinity of 2θ=30° to 35°. Specifically, when the ordered reflection peak intensity was higher than the background in that region, it was evaluated as having an ordered phase, and when the peak intensity was at or below the same level as the background, it was evaluated as not having an ordered phase.

[0092] As an example of the results of the XRD analysis performed in this embodiment, XRD diffraction patterns of the solution-treated and aged precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 are shown in FIG. 1. FIG. 1(a) shows an XRD diffraction profile for confirming spinodal decomposition, and FIG. 1(b) shows an XRD diffraction profile for confirming the ordered phase. Referring to FIG. 1(a), in this precious metal alloy, clear peaks identifiable as sideband peaks were confirmed on both sides of the peak corresponding to the {111} plane near 2θ = 40° to 42° and the peak corresponding to the {200} plane near 2θ = 46.5° to 48.5°. From this, it can be determined that spinodal decomposition occurs in this precious metal alloy. Furthermore, referring to Figure 1(b), a peak (ordered reflection peak) that is clearly higher than the background is observed in the region around 2θ = 31.5° to 34°, which indicates that ordering also occurs in this noble metal alloy.

[0093] [Study of material structure by TEM / STEM analysis] In order to confirm the occurrence of spinodal decomposition and ordering in the precious metal alloy produced in this embodiment, TEM / STEM analysis was performed. In the TEM / STEM analysis, a sample (φ0.6 mm) that had been subjected to solution treatment and aging treatment was subjected to resin embedding and polishing, and then a sample for TEM / STEM analysis was prepared using a focused ion beam (FIB), and the TEM / STEM analysis was performed.

[0094] FIG. 2 shows the results of mapping measurements of each constituent element (Au, Ni, Pd, Pt) of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 by STEM-EDS (Pt, Au, Pd: L line, Ni: K line). The electron incident direction is parallel to the

[001] direction. From FIG. 2, it can be seen that the material structure of the precious metal alloy of this embodiment has a modulated structure having two regions: a region where the Au concentration and Pd concentration are relatively high and a region where the Au concentration and Pd concentration are relatively low, and these regions are alternately connected. Since this modulated structure does not have a clear interface, it is presumed to be due to spinodal decomposition. Furthermore, comparison of the mapping image with the electron diffraction pattern confirmed that it has a modulated structure in the <010> direction.

[0095] FIG. 3 shows an electron diffraction pattern (001 zone axis incidence) obtained by TEM analysis of the precious metal alloy of Example 4-2. From FIG. 3, in this precious metal alloy, diffraction spots due to the ordered phase were confirmed in addition to the basic reflection due to the fcc structure. The diffraction spots of the ordered phase were identified as the L1 ordered phase structure from the appearance position, intensity, and interplanar spacing. ２ Furthermore, streaks in the <010> and <100> directions were observed at the 010 and 100 reciprocal lattice points, respectively, suggesting the existence of a modulated structure in the <010> direction. This result is consistent with the results of mapping measurements using STEM-EDS.

[0096] Similar results of the TEM / STEM analysis were obtained for other precious metal alloys. Therefore, it was confirmed that elemental analysis (composition mapping) and electron diffraction are also useful for determining spinodal decomposition and ordering in addition to XRD analysis (diffraction method).

[0097] [Confirmation of Segregation in Precious Metal Alloy After Solution Treatment] Figure 4 shows the results of STEM-EDS analysis of the vicinity of a high-angle grain boundary in the precious metal alloy (Au7.5-Ni41.25-Pd10-Pt41.25) of Comparative Example 2 after solution treatment. In Figure 4, the center of the photograph is the high-angle grain boundary, and while the concentrations of Pt and Ni are low near the grain boundary, the concentrations of Au and Pd are high near the grain boundary. This confirms that the Au-Ni-Pd-Pt-based precious metal alloy of this embodiment tends to have Au and Pd segregate near the grain boundary due to solution treatment.

[0098] Next, the presence or absence of segregation of B to grain boundaries was confirmed for the precious metal alloys of the examples to which B was added. In this embodiment, AES analysis (Auger electron spectroscopy) was performed to detect the light element B. The AES apparatus used was a PHI-700 model manufactured by ULVAC-PHI. In this embodiment, the precious metal alloy samples were impact fractured under ultra-high vacuum within the apparatus, and the fracture surfaces were observed, and AES analysis was performed on the grain boundary fracture surface and the intragranular fracture surface. To prepare the samples for AES analysis, a precious metal alloy cast to a diameter of 11 mm was subjected to the same homogenization and two-phase treatment as described above, followed by cold swaging to a diameter of 6 mm. Furthermore, the samples were subjected to intermediate annealing and cold groove rolling to a diameter of 3.75 mm, followed by solution treatment (1175°C for 10 minutes) to prepare evaluation samples. The temperature and time of this solution treatment are conditions for intentionally coarsening the crystal grains and facilitating the appearance of grain boundary fracture surfaces. The impact test was carried out in a cooled state at about -100°C. As a result, both intergranular and intragranular fracture surfaces were obtained on the fracture surface after impact fracture, and Auger spectra of both were measured.

[0099] As an example of the AES analysis results for the precious metal alloy of the example, Figure 5 shows the Auger spectrum of the grain boundary fracture surface and the intragranular fracture surface of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35). When comparing the spectra of the grain boundary fracture surface and the intragranular fracture surface, a peak near 185 eV was confirmed only in the grain boundary fracture surface. This peak near 185 eV is thought to be a peak derived from B, and therefore B is thought to be segregated at the grain boundaries of the recrystallized structure. This confirmed that B added to precious metal alloys tends to segregate at grain boundaries.

[0100] [Evaluation of ductility of recrystallized structure] The ductility of the wire rod after the solution treatment was evaluated. In the ductility evaluation test, a bending test (bending radius 1 mm, 180° bending) was performed on a wire rod sample with a diameter of 0.6 mm, and the bending angle until fracture was measured. The bending test was performed five times on the same sample, and the average value of the bending angle was evaluated as the bending value. Note that for samples that did not fracture even when bent to 180°, the bending value was taken as 180°.

[0101] [Hardness Measurement After Aging Treatment] The hardness of the wire samples after aging treatment was measured using a measuring device (HM-210 manufactured by Mitutoyo Corporation) at room temperature with a test load of 0.075 kgf. The measurement results are shown in Table 1. Hardness was measured at 15 random locations on each sample, and the average value was taken as the hardness value. For the measurement locations on each sample, multiple crystal grains were selected, and measurements were taken at the non-grain boundary reaction zone of each crystal grain, preferably near the center of the crystal grain. The hardness measurement results shown below include the aging treatment temperature and hardness at which the highest hardness was obtained among multiple aging treatment temperatures set in the range of 450°C to 600°C. However, for the precious metal alloys of the examples, even the lowest hardness was 500 Hv or more.

[0102] Tables 1 to 4 show the compositions of the precious metal alloys of the examples and comparative examples produced in this embodiment, as well as the results of ductility evaluation and hardness measurement. These tables also show the presence or absence of modulated structures (sideband peaks) and ordered phases (order peaks) determined by the XRD analysis described above. The alloy compositions listed in Tables 1 to 4 are analytical values ​​obtained for wire samples of each example. The alloy compositions were analyzed by ICP-OES analysis, and ICP-MS analysis was also used for the B concentration of samples expected to have trace amounts of B (0.05 atomic % or less). In these tables, in this embodiment, in order to confirm the effects of adding the essential additive element B, the optional additive element Cu, and the metal elements α and β, the precious metal alloys are grouped into examples and comparative examples, with the alloys being divided into groups based on the concentrations of the essential metal elements (Au, Ni, Pd, Pt) that are close to each other.

[0103]

[0104]

[0105]

[0106]

[0107] When examining the results shown in Tables 1 to 4, it should be noted that the Au—Ni—Pd—Pt alloys of each Example and Comparative Example, except for those that could not be solution-treated due to cracking during processing, all exhibited spinodal decomposition and / or ordering and exhibited hardness of 500 Hv or more. From this, it can be said that the Au—Ni—Pd—Pt alloys that the applicant focused on have suitable potential as high-hardness precious metal alloys.

[0108] Therefore, we will examine the effect of B, an essential additive element for improving ductility in the recrystallized structure after solution treatment, which is the subject of the present invention. From the table, it can be seen that the Au—Ni—Pd—Pt alloys of each Example to which B was added clearly exhibited increased bending angles compared to the Au—Ni—Pd—Pt alloys of Comparative Examples to which B was not added. For example, this point is clear when comparing Examples 1, 2, 3-1 to 3-5, and 4-1 to 4-4 in Table 1 and their corresponding Comparative Examples. The same is true when comparing the combinations of Examples and Comparative Examples in each table. Therefore, it can be said that the effect of adding B on improving ductility in the recrystallized structure after solution treatment has been confirmed.

[0109] However, the results in Table 1 confirm that the addition of B to Au—Ni—Pd—Pt alloys should be within a specified range. When the amount of B added is too small (0.001 atomic %), no effect of improving ductility is observed (Comparative Example 3). Moreover, when excessive B is added, the ductility of the two-phase structure, which is inherently good in workability, decreases, resulting in cracks during working (Comparative Example 11-2).

[0110] From the results of the investigation into the addition of B described above, it is believed that in the Au—Ni—Pd—Pt-based alloy of the present invention, the added B segregates at the grain boundaries, changing the grain boundary structure, and as a result, the grain boundary strength is improved, which suppresses grain boundary embrittlement and improves the ductility of the recrystallized structure.

[0111] The effects of additive elements other than B on the effectiveness of B, an essential additive element of the Au—Ni—Pd—Pt alloy of the present invention, were investigated. The results are shown in Comparative Examples 4-2 (additive element: C), 4-3 (additive element: C), 4-4 (additive element: Zr), and 4-5 (additive element: Ag) (see Table 1). When the evaluation results of the Au—Ni—Pd—Pt alloys of these comparative examples were examined, the addition of C and Zr did not improve ductility in the recrystallized structure (Comparative Examples 4-2 and 4-4). Furthermore, increasing the concentration of C caused grain boundary embrittlement, resulting in grain boundary cracking even in a two-phase structure, making processing difficult (Comparative Example 4-3). Furthermore, the addition of Ag, an element with a strong tendency for grain boundary segregation, was also attempted, but grain boundary embrittlement became apparent, resulting in grain boundary cracking even in a two-phase structure, making processing difficult (Comparative Example 3-5).

[0112] The effect of adding B to a recrystallized Au-Ni-Pd-Pt alloy on improving ductility has also been confirmed by room-temperature tensile testing. Figure 6 shows the results of room-temperature tensile tests conducted on the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) and the precious metal alloy of Comparative Example 4-1 (Au12.5-Ni37.5-Pd12.5-Pt37.5). The room-temperature tensile tests were conducted on solution-treated samples (φ0.6 mm × L50 mm) at a chuck distance of 25 mm and a crosshead speed of 10 mm / min. As can be seen from Figure 6, the addition of B significantly improved the elongation value, indicating a significant improvement in ductility.

[0113] Next, the effects of the optional added elements Cu, metal element α, and metal element β will be examined. Table 3 shows the results for precious metal alloys in which Cu was added to Au—Ni—Pd—Pt alloys. Table 3 confirms that spinodal decomposition and ordering also occur in Au—Ni—Pd—Pt alloys to which Cu was added. These precious metal alloys exhibit good hardness of 500 Hv or more after aging treatment.

[0114] The effect of adding B was also confirmed in the Au—Ni—Pd—Pt alloys containing Cu. The precious metal alloys without B added in the comparative examples had poor bendability and grain boundary embrittlement occurred, but it was confirmed that adding B as in each example improved bendability and suppressed grain boundary embrittlement.

[0115] Table 4 shows the effects of adding metal elements α (In, Sn, Sb) and metal elements β (Ti, Zr, Hf, Al). First, it was confirmed that none of these added elements inhibits the occurrence of spinodal decomposition and ordering in Au—Ni—Pd—Pt alloys. Furthermore, regardless of whether or not B is added, an increase in hardness is observed by adding at least one of these added elements. Depending on the composition of the precious metal alloy, this increase in hardness can reach 100 Hv or more.

[0116] It was also confirmed that the addition of B has the effect of improving ductility even in Au-Ni-Pd-Pt alloys to which the metal elements α and β have been added. Furthermore, a comparison between Example 24 and Example 26 in Table 3 shows that the metal elements α and β are also effective in increasing hardness in Au-Ni-Pd-Pt alloys containing the optional element Cu. The addition of either the metal element α or the metal element β may be either one or both, or multiple types of metal elements may be added.

[0117] As explained above, the present invention is a noble metal alloy with high hardness due to spinodal decomposition and / or ordering, and further has improved ductility in a recrystallized structure produced by solution treatment, etc. In the present invention, it is believed that, in response to grain boundary embrittlement due to grain boundary segregation of Au and Pd after solution treatment, the grain boundary structure is changed by the addition of B, which segregates B at the grain boundaries, and as a result, the grain boundary strength is improved, thereby suppressing grain boundary embrittlement.

[0118] The precious metal alloy according to the present invention is expected to be applicable to various applications requiring high hardness and high wear resistance, such as electrical and electronic materials such as probe pins and electrical contacts, medical instruments, and high-hardness coating members. In these applications, processing after heat treatment such as solution treatment is an essential step. The precious metal alloy according to the present invention exhibits good ductility after heat treatment such as solution treatment, making it effective for these applications. Furthermore, the precious metal alloy according to the present invention can be produced by aging a precious metal alloy obtained by additive manufacturing technology, metal powder injection molding technology, rapid solidification technology, laser heating technology, build-up welding technology, or coating technology such as sputtering, thermal spraying, or plating, and can be applied to the various applications described above.

Claims

1. An Au-Ni-Pd-Pt based precious metal alloy containing 1.5 atomic % or more and 47 atomic % or less of Au, 4 atomic % or more and 57 atomic % or less of Ni, 1 atomic % or more and 43.5 atomic % or less of Pd, and 6 atomic % or more and 58.5 atomic % or less of Pt, and further containing 0.003 atomic % or more and 8 atomic % or less of B.

2. The Au-Ni-Pd-Pt based precious metal alloy according to claim 1, further comprising more than 0.1 atomic % and not more than 27.5 atomic % of Cu.

3. An Au-Ni-Pd-Pt based precious metal alloy according to claim 1 or 2, further comprising 0.15 atomic % or more and 5 atomic % or less of a metal element α, said metal element α being at least any one of In, Sn and Sb.

4. An Au-Ni-Pd-Pt based precious metal alloy according to claim 1 or 2, further comprising 0.05 atomic % or more and 5 atomic % or less of a metal element β, said metal element β being at least any one of Al, Ti, Zr and Hf.

5. The Au-Ni-Pd-Pt based precious metal alloy according to claim 3, further comprising 0.05 atomic % or more and 5 atomic % or less of a metal element β, said metal element β being at least any one of Al, Ti, Zr and Hf.

6. A Au-Ni-Pd-Pt based precious metal alloy according to claim 1 or 2, which contains a modulated structure due to spinodal decomposition in its material structure.

7. A Au-Ni-Pd-Pt based precious metal alloy according to claim 1 or 2, which contains an ordered phase in its material structure.

8. The Au-Ni-Pd-Pt based precious metal alloy according to claim 6, which contains an ordered phase in its material structure.

9. The Au-Ni-Pd-Pt based precious metal alloy according to claim 1 or 2, which has a Vickers hardness of 500 Hv or more.

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