High-hardness Au-Ni-Pd-Pt-based precious metal alloy

Abstract

The present invention is an Au-Ni-Pd-Pt-based precious metal alloy comprising: 1.5 atomic percent or more and 47 atomic percent or less of Au; 4 atomic percent or more and 57 atomic percent or less of Ni; 1 atomic percent or more and 43.5 atomic percent or less of Pd; and 6 atomic percent or more and 58.5 atomic percent or less of Pt, wherein the Au-Ni-Pd-Pt-based precious metal alloy further comprises 0.003 atomic percent or more and 8 atomic percent or less of B. The present precious metal alloy according to the invention exhibits increased hardness due to a modulated texture resulting from spinodal demixing and / or an ordered phase generated by order formation. In addition to the increased hardness, the invention exhibits improved ductility in a recrystallized texture due to the addition of B. The precious metal alloy of the present invention may optionally contain Cu, an additive metal α (In, Sn, Sb) and an additive metal β (Al, Ti, Zr, Hf).

Description

AREA OF INVENTION

[0001] The present invention relates to a precious metal alloy with high hardness, in which Pt, Au and Pd are essential constituent metals, taking into account workability, in particular an Au-Ni-Pd-Pt-based precious metal alloy which has achieved increased hardness by spinodal demixing and / or order formation and which has improved workability thanks to increased ductility of a recrystallized texture, which has been achieved by solution annealing and the like. DESCRIPTION OF THE RELATED AREA

[0002] Precious metals such as platinum (Pt) and gold (Au) are not only excellent in terms of chemical stability and corrosion resistance, but also advantageous in terms of electrical properties such as conductivity. Therefore, precious metals and their alloys are used in various fields, including electrical engineering / electronics and medicine. Examples of precious metal alloy applications in electrical engineering / electronics include probes used for testing semiconductors, for example, in front-end and back-end processes, and electrical contacts (sliding and make-and-break contacts) for motor brushes, relays, switches, and the like. Their use in medicine has recently gained attention, and precious metal alloys are used as constituent materials in various medical instruments.Examples of such medical instruments include various types of medical devices, such as embolization coils and clips, guide wires, stents, and catheters. These medical instruments are intended to come into direct contact with and be embedded in the human body and therefore must exhibit biocompatibility and chemical stability. Such medical instruments must also be radiopaque, considering their use in surgery and diagnostics with X-rays. Precious metal alloys are also advantageous with regard to biocompatibility and radiopaqueness.

[0003] Precious metal alloys intended for the various applications described above must exhibit enhanced mechanical properties such as hardness and strength. For example, styluses are intended to repeatedly contact their counterparts over extended periods and therefore must possess wear resistance. In particular, it is necessary to develop styluses with higher hardness to accommodate the high integration of modern devices and the high performance of modern motors. Medical instruments such as guide wires and embolization coils, which are moved and then embedded in pulsating / beating vessels, must possess mechanical properties such as hardness and spring characteristics to ensure the reliable operation of such instruments.

[0004] Known methods for increasing the hardness of metallic materials, including precious metal alloys, are material strengthening techniques such as strain strengthening (dislocation strengthening), solid solution strengthening, and precipitation strengthening (dispersion strengthening), and these methods are applied either individually or in combination. As examples of hardness enhancement in precious metal alloys used for the aforementioned styluses, contact materials, medical equipment, and the like, a Pt-Ni alloy described in Patent Document 1 and a Pt-W alloy in Patent Document 2 achieve increased hardness through strain strengthening at a high finishing rate, in addition to solid solution strengthening by alloying Pt with Ni, W, or the like.In patent document 3 (Ag-Pd-Cu-based alloy) and patent document 4 (Pt-Cr-Ni-based alloy), a precious metal alloy with high hardness is achieved not only by solid solution strengthening and precipitation hardening with additional elements, but also by deformation strengthening, whereby the processing rate is adjusted.

[0005] In contrast, the present applicant has developed a precious metal alloy with high hardness, which is achieved by spinodal demixing and / or ordering (patent document 5) as a material strengthening process different from the above-mentioned known techniques.

[0006] Spinodal demixing corresponds to a type of phase separation in a material texture and is a phenomenon in which demixing progresses due to a continuous increase in concentration variations. A material texture generated by spinodal demixing due to concentration variations exhibits a very fine periodic structure ranging from a few nanometers to several tens of nanometers, which is referred to as a modulated texture. A modulated texture resulting from spinodal demixing varies periodically with respect to the concentration of a dissolved atom in a crystal as a function of position and also changes periodically with respect to the lattice constant. Thus, a periodic internal stress field is generated at a slip surface, and this field interacts with dislocation.The solidification mechanism of spinodal demixing is similar to excretory solidification by nucleation / growth, but differs in that internal stress fields created by concentration modulation are used as obstacles to dislocation movement instead of precipitates.

[0007] Order formation is a phenomenon in which the arrangement of constituent elements in an alloy is ordered to create an ordered phase with a predetermined structure. The ordered phases generated by order formation contribute to an increase in the hardness of an alloy due to each of the following factors: (i) an increase in the Burgers vector of the dislocation, (ii) the potential appearance of an antiphase boundary within the ordered phase, and (iii) volume changes associated with order formation, which distort the lattice both within and outside the ordered phase, resulting in the hindrance of dislocation movement. Order formation can occur in conjunction with the spinodal demixing described above, or it can occur independently.

[0008] Spinodal segregation and ordering are phenomena that, due to unique material textures such as modulated textures and ordered phases, can cause a significant increase in hardness. Solid solution strengthening and precipitation strengthening, i.e., well-known material strengthening processes, have limitations regarding the degree of hardening. Spinodal segregation and ordering can be expected to produce a greater degree of hardening compared to these conventional strengthening methods. Furthermore, although strain strengthening can be expected to increase hardness with a higher degree of processing, concerns exist regarding material embrittlement due to the hardening process. Spinodal segregation and ordering are considered useful procedures for increasing hardness without causing material embrittlement.

[0009] The precious metal alloy proposed by the present applicant (Patent Document 5) incorporates an Au-Ni-Pd-Pt-based alloy as a quaternary or multi-component alloy to effectively exhibit the work-hardening effects of spinodal segregation and ordering. Binary alloys such as Pt-Au and Pt-Ni alloys are conventionally known as precious metal alloys capable of exhibiting spinodal segregation and ordering, but the degree of hardening of these alloys is not particularly high. The spinodal segregation and ordering in these binary alloys are not strong enough to outperform conventional work-hardening methods such as strain hardening and precipitation hardening.In the Au-Ni-Pd-Pt-based alloy proposed by the present applicant, the use of a multi-component alloy, in contrast to the binary alloys mentioned above with optimization of the composition range of each constituent element, allows the hardening effects of spinodal demixing and / or order formation to be effectively demonstrated. In the Au-Ni-Pd-Pt-based alloy, spinodal demixing and order formation can occur by first forming a supersaturated solid solution in a recrystallized texture state from an alloy with a suitable composition through solution annealing or the like, and then performing a subsequent aging treatment, thus achieving the desired hardness. State of the art document / Patent document Patent document 1 Japanese patent application, publication number 2005-233967 Patent document 2 Japanese Patent No. 6997354 Patent document 3 Japanese patent application Disclosure number 2012-242184 Patent document 4 Japanese Patent No. 6372952 Patent document 5 International Publication No. WO2023 / 063156 SUMMARY OF THE INVENTIONAL PROBLEM

[0010] When supplying precious metal alloys for styluses, sliding contacts, medical equipment, and the like, it is necessary to process the alloys plastically into the shapes and dimensions required for the specific applications. Regarding the processability of the aforementioned Au-Ni-Pd-Pt-based alloy proposed by the present applicant, this precious metal alloy possesses sufficient ductility even at room temperature, both in a cast textured state and in a two-phase separated textured state, which occurs after heat treatment at approximately 700 to 900 °C, and thus exhibits good processability.Therefore, to process this Au-Ni-Pd-Pt-based alloy into styluses, sliding contacts, medical equipment, and the like, primary processing such as rolling and wire drawing is carried out on the alloy in a state (either the cast or the two-phase separated texture) with good workability to produce raw materials such as wire bars. It is assumed that after this primary processing, the raw materials undergo a solution annealing treatment, followed by secondary and final processing such as dressing, straightening, bending, and winding, and that the materials are subsequently subjected to an aging treatment to harden them.

[0011] However, investigations carried out by the present inventors have confirmed that the Au-Ni-Pd-Pt-based alloy proposed by the present applicant exhibits a disadvantage with regard to ductility in the post-solution annealing state. This reduction in ductility can affect processing and may diminish workability during the secondary processing and similar operations described above.

[0012] In the Au-Ni-Pd-Pt-based alloy proposed by the present applicant, the alloy is hardened by spinodal segregation and / or ordering in an aging treatment, and to maximize this effect, a solution annealing treatment performed prior to the aging treatment is an essential step. Accordingly, assuming that the processing procedure described above is applied to the Au-Ni-Pd-Pt-based alloy proposed by the present applicant, an improvement in ductility after the solution annealing treatment is considered a critical point.

[0013] The present invention is based on the above circumstances and provides a precious metal alloy with high hardness comprising the Au-Ni-Pd-Pt-based alloy proposed by the present applicant, which, in a state after solution annealing and the like, is improved with respect to ductility while exhibiting the high hardness achieved by spinodal demixing and / or ordering. SOLUTION TO PROBLEM

[0014] The Au-Ni-Pd-Pt-based alloy proposed by the present applicant (patent document 5) is an unprecedented technology in that the effects of spinodal demixing and / or order formation are actively utilized for a precious metal alloy. Therefore, the phenomenon of reduced ductility after solution annealing described above is itself unknown. To address this issue, the inventors decided to clarify the cause of the reduction in ductility of the Au-Ni-Pd-Pt-based alloy after solution annealing.

[0015] As a result, the present inventors have found that in an Au-Ni-Pd-Pt-based alloy capable of exhibiting spinodal demixing and order formation, grain boundary segregation of Au and / or Pd can occur in a recrystallized texture, i.e., a material texture produced by solution annealing. The inventors therefore hypothesize that this grain boundary segregation causes grain boundary embrittlement in the recrystallized texture, which can be a cause of reduced ductility. This type of recrystallized texture is a material texture found in all or part of the Au-Ni-Pd-Pt-based alloy after it has undergone solution annealing and similar treatments.

[0016] Assuming that the reduced ductility following solution annealing is caused by grain boundary embrittlement due to the segregation of Au and / or Pd in ​​the recrystallized texture, the prerequisite for a countermeasure is to improve ductility without inhibiting the hardening capacity of spinodal segregation and order formation. In other words, even to suppress the segregation of Au and Pd, it is not a preferred approach to investigate changing the composition of Au, Pd, and other constituent elements to a compositional range where neither spinodal segregation nor order formation can occur. Therefore, the inventors sought to find an element for the Au-Ni-Pd-Pt-based alloy that would segregate at the grain boundaries and strengthen these grain boundaries without inhibiting spinodal segregation or order formation.As a result, it was discovered that B (boron) meets the above conditions, and thus the present invention was conceived.

[0017] Specifically, the present invention relates to an Au-Ni-Pd-Pt-based precious metal alloy containing 1.5 atomic% or more and 47 atomic% or less Au, 4 atomic% or more and 57 atomic% or less Ni, 1 atomic% or more and 43.5 atomic% or less Pd and 6 atomic% or more and 58.5 atomic% or less Pt, and containing 0.003 atomic% or more and 8 atomic% or less B.

[0018] The precious metal alloy of the present invention contains Au, Ni, Pd, Pt and B as essential constituent elements, and it is also acceptable even if it contains some additional elements, and can contain more than 0.1 atomic % and 27.5 atomic % or less Cu.

[0019] The Au-Ni-Pd-Pt-based precious metal alloy of the present invention can further comprise, as optional additive elements, 0.15 atomic percent or more and 5 atomic percent or less of a metal element α and / or 0.05 atomic percent or more and 5 atomic percent or less of a metal element β. Here, the metal element α is at least one metal element selected from In, Sn, and Sb, and the metal element β is at least one metal element selected from Al, Ti, Zr, and Hf.

[0020] As described above, the Au-Ni-Pd-Pt-based precious metal alloy of the present invention contains a modulated texture through spinodal demixing and / or an ordered phase.

[0021] The Au-Ni-Pd-Pt-based precious metal alloy of the present invention exhibits increased hardness due to spinodal demixing and / or ordering. This Au-Ni-Pd-Pt-based precious metal alloy with thus increased hardness has a Vickers hardness of 500 Hv or more. ADVANTAGEOUS EFFECTS OF THE INVENTION

[0022] As described above, the Au-Ni-Pd-Pt-based precious metal alloy of the present invention exhibits improved ductility in the recrystallized texture after solution annealing. Similar to the Au-Ni-Pd-Pt-based precious metal proposed by the present applicant, the present invention allows for increased hardness due to a modulated texture resulting from spinodal demixing and / or an ordered phase resulting from order formation, both of which are achieved through appropriate heat treatment, such as aging. The present invention can ensure processability into any shape according to various uses, even in a recrystallized texture state after solution annealing, and suitable hardness can be obtained through subsequent aging. BRIEF DESCRIPTION OF THE DRAWINGS Fig. Figure 1 is a diagram showing the results of X-ray diffraction (XRD) of a solution-annealed material and an aged material of a precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35); Fig. Figure 2 is a STEM-EDS distribution image of a sample of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35), which has been subjected to an aging treatment after a solution annealing treatment; Fig. Figure 3 is an electron diffraction pattern of the sample of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35), which has been subjected to an aging treatment after a solution annealing treatment; Fig. 4 are photographs showing results of the STEM-EDS analysis of a precious metal alloy of comparison example 2 (Au7.5-Ni41.25-Pd10-Pt41.25) and illustrating each element distribution near a high-angle grain boundary; Fig. Figure 5 is a diagram showing Auger spectra of the intergranular and transgranular fracture surfaces after impact fracture of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35); and Fig. Figure 6 is a diagram showing the results of tensile tests at room temperature of the precious metal alloys of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) and the precious metal alloy of the comparison example 4-1 (Au12.5-Ni37.5-Pd12.5-Pt37.5). DETAILED DESCRIPTION OF PREFERRED EXECUTION FORMS

[0023] Embodiments of the present invention are described below. As described above, the precious metal alloy of the present invention relates to an Au-Ni-Pd-Pt-based alloy resulting from the addition of an additive element B to an Au-Ni-Pd-Pt alloy. The configuration of the precious metal alloy of the present invention and a method for producing the precious metal alloy according to the invention are now described. (A) Configuration of the present precious metal alloy according to the invention (A-1) Respective constituent elements and composition ranges of the present precious metal alloy according to the invention

[0024] The precious metal alloy of the present invention contains Au, Ni, Pd, and Pt as essential constituent elements and B as an essential additive element. The following description clarifies the essential constituent elements and their composition ranges. Optional additive elements (Cu, a metal element α, and a metal element β), the addition of which to the present precious metal alloy according to the invention is acceptable in addition to B, as well as their additive concentrations, are then described. (1) Essential constituent elements (Au, Ni, Pd and Pt)

[0025] The precious metal alloy of the present invention contains these essential elements and therefore exhibits at least one of spinodal segregation and order formation, and it possesses a hardness equal to or greater than that of precious metal alloys achieved by conventional strengthening mechanisms (solid solution strengthening, precipitation strengthening and strain strengthening).

[0026] As described in patent document 5, which discloses the Au-Ni-Pd-Pt-based alloy proposed by the present applicant, a modulated texture formed by spinodal demixing varies periodically with respect to concentration, forming a perimeter internal stress field and thus contributing to an increase in hardness. The resistance (critical shear stress) to dislocation movement in such a periodic internal stress field is expressed by the following expression, and it is assumed that the lattice strain (ε), the elasticity coefficient (Y), and the concentration modulation amplitude (A) are controlling factors (examples of specific reference documents include Masaharu Kato, Introduction to the Theory of Dislocations (published in August 1999, publication: Shokado)). τ=YA|ε|6 τ Critical shear stress Y elasticity coefficient A Concentration modulation amplitude ε Lattice strain

[0027] Considering the above expression, the constituent elements of a noble metal alloy that effectively exhibits material strengthening by spinodal demixing should preferably meet the following three requirements: (1) The constituent elements should include a metal with a high elasticity coefficient; (2) the enthalpy of mixing between the constituent elements should be high, and phase separation should strongly tend to occur in a low temperature range; and (3) there should be a large difference in the lattice constant between the constituent elements, and the lattice strain (ε) should also be large.

[0028] Regarding the occurrence of order formation, the combinations of Pt and Ni as well as Au and Pd are metal pairs that can contribute to the generation of an ordered phase through order formation.

[0029] The essential constituent elements of the precious metal alloy of the present invention, namely Au, Ni, Pd, and Pt, are a combination of metals that can meet these requirements, and when a composition range as described below is used, the work-hardening capacity through spinodal demixing is effective. The specific effects and composition ranges of these essential constituent elements are explained below. Au

[0030] Au is an essential element for the occurrence of spinodal segregation in the alloy system of the present invention. Spinodal segregation does not occur at excessively high or low Au concentrations, and there is a concentration range for Au that is necessary for its occurrence. If the Au concentration is outside this optimal range, typical nucleation / growth readily occurs, and a suitable increase in hardness cannot be achieved. Furthermore, Au and Pd are metals capable of forming an ordered phase and also contribute to an increase in hardness through order formation.

[0031] The Au concentration in the Au-Ni-Pd-Pt-based alloy of the present invention is 1.5 atomic % or more and 47 atomic % or less, and is preferably 4 atomic % or more and 30 atomic % or less, more preferably 6 atomic % or more and 24 atomic % or less, and further preferably 7 atomic % or more and 20 atomic % or less. Ni

[0032] Ni acts as a work-hardening factor in the occurrence of spinodal segregation in the precious metal alloy. Ni has a higher coefficient of elasticity than Au, Pt, and Pd. Furthermore, compared to Au, Pt, and Pd, Ni has a smaller lattice constant and therefore has a large effect on increasing the lattice strain ε. Accordingly, Ni acts in such a way that it increases the work-hardening capacity due to spinodal segregation, as can be seen from expression 1. In addition, Ni and Pt are metals capable of forming an ordered phase and also act in such a way that they contribute to an increase in hardness through ordering.

[0033] Ni is a congener of Pt and Pd and resembles them in electronic structure, thus enabling the formation of an alloy without any loss of corrosion and oxidation resistance comparable to a precious metal. Consequently, Ni also has the secondary effect of reducing the overall price of the precious metal alloy.

[0034] The Ni concentration in the Au-Ni-Pd-Pt-based alloy of the present invention is 4 atomic % or more and 57 atomic % or less, and is preferably 10 atomic % or more and 52.5 atomic % or less, more preferably 12.5 atomic % or more and 50 atomic % or less, and further preferably 15 atomic % or more and 47.5 atomic % or less. Pd

[0035] Pd has the effects of extending the solid-state solubility limit of each element constituting the precious metal alloy, widening the concentration range that allows spinodal segregation to occur in the alloy, and promoting spinodal segregation. Thanks to these effects, Pd increases the degree of hardening due to spinodal segregation. Furthermore, gold and Pd are metals capable of forming an ordered phase and thus contribute to increased hardness through ordering. However, excessive addition of Pd excessively reduces the temperature of spinodal segregation, and consequently, spinodal segregation tends to be inhibited. Moreover, excessive addition of Pd can also suppress ordering, leading to a reduction in the overall hardening of the alloy system.Accordingly, Pd also has an optimal concentration range as described above in order to optimize the degree of hardening of the precious metal alloy.

[0036] The Pd concentration in the Au-Ni-Pd-Pt-based alloy of the present invention is 1 atomic % or more and 43.5 atomic % or less, and is preferably 3 atomic % or more and 37.5 atomic % or less, more preferably 4 atomic % or more and 32.5 atomic % or less, and further preferably 5 atomic % or more and 27.5 atomic % or less. Pt

[0037] Platinum (Pt) is also an essential element for spinodal segregation in the alloy system of the present invention. Spinodal segregation does not occur at excessively high or low Pt concentrations, and there is a concentration range of Pt necessary for its occurrence. Pt can be combined with nickel to form an ordered phase, which contributes to an increase in hardness. Furthermore, Pt exhibits a relatively high coefficient of elasticity, and it can therefore be expected to be an element that allows for an increase in the degree of hardening of the alloy in which spinodal segregation occurs.

[0038] The Pt concentration in the Au-Ni-Pd-Pt-based alloy of the present invention is 6 atomic % or more and 58.5 atomic % or less, and is preferably 12.5 atomic % or more and 52.5 atomic % or less, more preferably 15 atomic % or more and 50 atomic % or less, and further preferably 17.5 atomic % or more and 47.5 atomic % or less. (2) Essential additional element (B)

[0039] B, i.e., an essential additional element in the Au-Ni-Pd-Pt-based alloy of the present invention, is an element which strengthens grain boundaries of the precious metal alloy by segregation at the crystal grain boundaries of a recrystallized texture formed by solution annealing or the like. This suppresses a reduction in ductility or embrittlement that can occur due to grain boundary segregation of Au and Pd in ​​a recrystallized section of the precious metal alloy after solution annealing, and confers suitable workability on the precious metal alloy containing the recrystallized texture.

[0040] The B concentration in the precious metal alloy of the present invention is 0.003 atomic percent or more and 8 atomic percent or less. If the concentration is less than 0.003 atomic percent, the degree of segregation at the grain boundaries of the recrystallized texture is so low that the grain boundary strengthening effect mentioned above cannot be sufficiently achieved. Conversely, a concentration of more than 8 atomic percent is not preferred because the ductility tends to decrease not only in the recrystallized texture state but also in the cast texture and two-phase texture states. The B concentration is preferably 0.0075 atomic percent or more. In addition, the B concentration can be appropriately set, taking into account the overall alloy composition of the Au-Ni-Pd-Pt-based alloy, 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, and the like. (3) Optional additional elements (Cu, metal element α and metal element β)

[0041] The Au-Ni-Pd-Pt-based alloy of the present invention uses a precious metal alloy comprising the essential constituent elements described above as a base alloy and may additionally contain optional additive elements. Examples of the optional additive elements are Cu, metal element α, and metal element β. Cu

[0042] Copper (Cu) has the effect of widening the solid-state solubility limit between the constituent elements of the precious metal alloy, thereby extending the concentration range for spinodal demixing. In this respect, Cu has a similar effect to Pd and can therefore be considered an additive element. However, although some of its effects are similar to Pd, Cu is an optional additive element and not an essential constituent element like Pd. This is because, while Cu widens the concentration range for spinodal demixing, it does not increase the degree of hardening through spinodal demixing.As an optional additional element, Cu, in addition to the effects described above, also has the effect of improving machinability and is also a less expensive metal than precious metals, and therefore has the secondary effect of reducing the overall price of the resulting precious metal alloy.

[0043] When copper is added as the optional additive element to achieve the effects mentioned above, the concentration is preferably greater than 0.1 atomic percent and less than 27.5 atomic percent. Excessive addition of copper can reduce the spinodal demixing temperature, leading to an inhibition of order formation and potentially decreasing the overall hardening of the alloy system. Additionally, excessive addition of copper can be a factor in reducing the corrosion resistance of the alloy. When copper is added, the concentration is more preferably greater than 2.5 atomic percent and less than 22.5 atomic percent, and more preferably greater than 5 atomic percent and less than 18.5 atomic percent. Metal element α and metal element β

[0044] The Au-Ni-Pd-Pt-based alloy of the present invention can, in addition to the essential constituent elements and Cu described above, contain the metal elements α and β as optional additive metal elements. The metal elements α and β have the effect of increasing the hardness of the precious metal alloy after aging treatment. While the reason for the increase in hardness of the precious metal alloy due to the addition of the metal elements α and β is not clear, the inventors of the present invention assume that it is related to: increased strength of an alloy matrix due to an increase in lattice strain; refinement of the crystal grains of the alloy after aging treatment; formation of precipitates containing the additive elements; and the like. Furthermore, the addition of the metal elements α and β does not inhibit the formation of a material texture that includes a modulated texture and / or an ordered phase.It can contain either one or both of the metallic elements α and β.

[0045] The metallic element α is at least one selected from In, Sn, and Sb. Investigations carried out by the present inventors have demonstrated an increase in the hardness of the precious metal alloy of the present invention when using these metallic elements. It is preferred to add one or more of In, Sn, and Sb as the metallic element α. The composition range for at least one of In, Sn, and Sb used as the metallic element α is 0.15 atomic percent or more and 5 atomic percent or less. An added amount of less than 0.15 atomic percent is unlikely to contribute to an increase in the hardness of the precious metal alloy, while an amount of more than 5 atomic percent causes a noticeable reduction in workability. The composition range for the metallic element α is preferably 0.3 atomic percent or more and 3 atomic percent or less, and more preferably 0.5 atomic percent or more and 1.5 atomic percent or less.

[0046] The metallic element β is at least one of Ti, Zr, Hf, and Al. Investigations carried out by the inventors revealed an improvement in the crystal grain size and hardness of the precious metal alloy of the present invention when using these metallic elements. It is preferred to add one or more of the metallic elements Ti, Zr, Hf, and Al as the metallic element β. The composition range for at least one element from Ti, Zr, Hf, and Al used as the metallic element β is 0.05 atomic percent or more and 5 atomic percent or less. An addition of less than 0.05 atomic percent is unlikely to contribute to a refinement of the crystal grain size or an increase in the hardness of the precious metal alloy, while an amount of more than 5 atomic percent causes a noticeable reduction in 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.

[0047] The precious metal alloy of the present invention is a precious metal alloy containing Au, Ni, Pd, Pt, and B as its essential elements within the composition ranges described above, and is preferably a precious metal alloy comprising Au, Ni, Pd, Pt, and B within the ranges described above. However, the precious metal alloy of the present invention may also contain any unavoidable impurities. An unavoidable impurity is an unavoidable component that is included in impurities in a raw material or is included as a result of a production step or the like. Specific examples of unavoidable impurities include Rh, Ir, Fe, Y, Zn, Co, Si, Cu, Th, H, and rare earth elements. Such unavoidable impurities are introduced from a raw material and apparatus and the like during melting and casting.The content of these unavoidable impurities is preferably within a range that does not impair the properties of the precious metal alloy of the present invention; the content per element is preferably less than 0.05 atomic percent, and the total content is preferably 0.5 atomic percent or less, and particularly preferably less than 0.1 atomic percent. When the unavoidable impurity is included in the precious metal alloy, it is difficult to clearly distinguish whether the unavoidable impurity is an inherent component or a deliberately added component. In the present invention, as long as the component does not significantly alter the properties of the precious metal alloy, it is assumed to be an unavoidable impurity, without any distinction as to the intention behind its introduction.This definition of unavoidable impurities also applies when the above-mentioned additional elements Cu and the metallic elements α and β are present.

[0048] Spinodal demixing and order formation of the precious metal alloy of the present invention occur in a solution annealing treatment for quenching a solid solution alloy with the composition range described above, after being held at a high temperature, and a subsequent aging treatment step. The precious metal alloy of the present invention has a region that allows a complete solid solution extending well into a high-temperature range and also has a miscibility gap in a low-temperature range. Therefore, it is assumed that by quenching the precious metal alloy with the aforementioned composition range after a solution annealing treatment at a high temperature, a supersaturated solid solution is formed, and then spinodal demixing and order formation can occur through the subsequent aging treatment.Here, it is also effective to apply a CALPHAD method (method for calculating phase diagrams) with respect to the thermodynamic behavior (transition point, phase equilibrium, limit of solubility in the solid state, melting point, and the like) of the precious metal alloy of the present invention. Calculation by a CALPHAD method is preferably carried out using commercially available thermodynamic calculation software (e.g., Thermo-Calc and precious metal alloy databases (e.g., TCNOBL1)). (A-2) Material texture of the present precious metal alloy according to the invention

[0049] The precious metal alloy of the present invention achieves an increase in hardness through spinodal demixing and / or ordering. Therefore, the material texture of the precious metal alloy of the present invention can comprise a finely modulated texture through spinodal demixing and / or an ordered phase through ordering.

[0050] In the precious metal alloy of the present invention, the modulated texture generated by spinodal demixing is a material texture whose composition varies with a modulation cycle on the order of a few nanometers to several tens of nanometers. Specifically, the modulated texture in the present invention is configured with two regions, one of which has relatively high Au and Pd concentrations (the region in which Pt and Ni concentrations are relatively low) and the other of which has relatively low Au and Pd concentrations (the region in which Pt and Ni concentrations are relatively high).

[0051] The composition and configuration of the ordered phase in the present invention are not necessarily fully known. However, it is assumed that a phase with the same or a similar crystal structure to an ordered phase that can be produced in Pt-Ni-based or Au-Pd-based alloys, which are known as precious metal alloys in which ordering can occur, is also produced in the precious metal alloy of the present invention. Accordingly, it is assumed that the ordered phase in the present invention has an L10-type or L12-type structure, or a crystal structure similar to these.

[0052] The material texture of the precious metal alloy of the present invention can be confirmed by X-ray diffraction (XRD), a transmission electron microscope (TEM), an electron diffraction pattern using TEM, a scanning transmission electron microscope (STEM), a TEM / STEM-EDS distribution image, a TEM / STEM-EELS distribution image, a 3D atom probe or the like.

[0053] The modulated structure of the precious metal alloy of the present invention can be confirmed by one or both of a diffraction method and elemental analysis (element mapping).

[0054] Diffraction methods include a procedure by confirming an X-ray diffraction pattern obtained by X-ray diffraction (XRD) or an electron diffraction pattern obtained by TEM. When using the diffraction method for confirmation, if a modulated structure is present, a broad peak, referred to as a sideband peak (satellite peak) or fringe, is observed on at least one side (preferably both sides) of either a main peak or an ordered reflection peak. The presence or absence of such a sideband peak can be used to determine whether a modulated texture is present due to spinodal demixing.For example, in a diffraction pattern obtained by X-ray diffraction, the crystal structure of a matrix of the precious metal alloy of the present invention is a face-centered cubic (fcc) lattice structure, and Miller indices {111}-plane, {200}-plane, {220}-plane, {311}-plane, and the like appear as main peaks. A sideband peak, which is associated with a texture by spinodal demixing, appears on both sides or on one side of at least one of the main peaks described above. The reason why such a sideband peak appears on only one side of such a main peak is assumed to be that separation from the main peaks is difficult.

[0055] Alternatively, an example of a confirmation procedure using element mapping includes a method in which differences in atomic concentration density are visually verified using STEM-EDS / EELS or a 3D atom probe.

[0056] To confirm the presence of an ordered phase, a procedure is typically used to check for the presence or absence of an ordered reflection peak in an X-ray diffraction pattern or an electron diffraction pattern obtained with a TEM. For example, when observing an ordered reflection peak in an X-ray diffraction pattern, an ordered reflection peak appears at 2θ = approximately 22.5 to 27.5°, approximately 30° to 35°, and so on in the case of a θ-2θ measurement, where CuKα radiation is used as an X-ray source. In an electron diffraction pattern, diffraction spots such as 100, 110, and 120 appear, which do not appear in an fcc structure. (A-3) Hardness of the present precious metal alloy according to the invention

[0057] The precious metal alloy of the present invention exhibits the above-mentioned compositional range and thus achieves hardening through spinodal segregation and / or order formation. The precious metal alloy of the present invention can stably exhibit a hardness of 500 Hv or more in the sense of Vickers hardness. The Vickers hardness of the precious metal alloy of the present invention is preferably 540 Hv or more, more preferably 590 Hv or more, and further preferably 640 Hv or more. The precious metal alloy of the present invention can achieve the high hardness with the above-mentioned Vickers hardness solely through heat treatment without any use of strain hardening, namely without the occurrence of material embrittlement due to dislocation strain.

[0058] The upper limit of the hardness of the precious metal alloy of the present invention is not to be particularly limited, and the upper limit is preferably 850 Hv or less. If the value is greater than 850 Hv, fractures and chipping may occur during use. The Vickers hardness described above is a value at room temperature. The Vickers hardness can be measured with a known Vickers hardness tester. The test load is preferably 0.025 kgf or more and 0.5 kgf or less, more preferably 0.075 kgf.

[0059] The shape and form of the precious metal alloy of the present invention are not specifically limited. For the aforementioned medical tools, styluses, and the like, the present invention can generally be used as an alloy in the form of a mass (a lump) that has undergone appropriate processing. In the present invention, the ductility is improved to allow processing into shapes and dimensions suitable for these uses. Alternatively, the present invention can also be formed in a layered / film form on a suitable base material / substrate. (B) Method for producing the present precious metal alloy according to the invention

[0060] Next, a method for producing a precious metal alloy according to the present invention is described. As described above, the precious metal alloy of the present invention can be provided in various forms and shapes. The following description provides details of a method for producing a precious metal alloy in bulk form, which is frequently used, and also mentions a method for producing a precious metal alloy in layered / film form.

[0061] In the present invention, hardness is increased by spinodal segregation and / or ordering in addition to the selection of the constituent elements (essential constituent elements (Au, Ni, Pd, Pt, and B)) and optional additive metallic elements (Cu and the metallic elements α and β) of the precious metal alloy and the optimization of its composition ranges. The precious metal alloy of the present invention is produced by carrying out an appropriate heat treatment step in addition to producing an alloy mass (block) with the specified composition ranges. The appropriate heat treatment step is a heat treatment step combining a solution annealing treatment and an aging treatment, and these are carried out to result in the progression of spinodal segregation and / or ordering to increase the hardness, thus obtaining the precious metal alloy of the present invention.

[0062] The following describes the process for producing the precious metal alloy of the present invention, specifying the respective heat treatment steps of the solution annealing and aging treatments. In the present invention, the temperature, for example, the heating temperature of various heat treatments described below, is that of the precious metal alloy to be treated, unless otherwise specified. (B-1) Provisioning step (production of the precious metal alloy)

[0063] First, a precious metal alloy block is provided prior to the process, serving as a precursor to the precious metal alloy of the present invention. This precursor block can be produced by a conventional melting and casting process. The respective raw metal materials of Au, Ni, Pd, Pt, B, and optional elements of Cu, and the metal elements α and β described above, are appropriately weighed or the like to allow for adjustment to the composition described above, and the resulting mixture is melted / cast to produce an alloy block. Here, alloys such as precious metal alloys that do not contain Cu and the metal elements α and β (an Au-Ni-Pd-Pt alloy) or binary alloys of an Au-Pd alloy, a Pt-Ni alloy, and the like can be appropriately combined as parent alloys to be melted.The melting and casting of such a precious metal alloy can be carried out by a known procedure such as arc melting, high-frequency melting, vacuum melting, continuous casting or a liquid quenching process.

[0064] Such a precious metal alloy can also be provided by any process other than melting and casting, such as a powder metallurgy process. In the powder metallurgy process, a precious metal alloy powder (for example, a precious metal alloy powder produced by atomization), which is modified to have the above composition, is sintered, and then an alloy ingot is provided for heat treatment. Alternatively, such a precious metal alloy powder, modified to have the above composition, can also be used to produce an ingot with a near-net-shape form by a known metal injection molding process or an additive manufacturing process.Furthermore, a precious metal alloy layer with the above composition can also be formed on any parent material by a known alloying process such as plating, sputtering or thermal spraying. (B-2) Solution annealing

[0065] A supersaturated solid solution of such a precious metal alloy, provided as above, is formed by a solution annealing treatment. The solution annealing treatment consists of heating such a precious metal alloy to a high temperature to obtain a texture with a high solid solution concentration, and then quenching the texture to form a supersaturated solid solution. The heating temperature of the solution annealing treatment is preferably a temperature of (Tm - 500 °C) or higher and Tm or lower, assuming that the melting point (solidus line) of such a precious metal alloy is Tm (°C). A temperature of less than (Tm - 500 °C) results in low solid solubility of each element and is insufficient for the formation of the supersaturated solid solution, and a temperature greater than Tm is not preferred because melting of the material begins near a grain boundary.The holding time during heating is preferably in the range of 0.0001 hours or more up to 168 hours. A holding time of less than 0.0001 hours is not preferred from a productivity point of view, since the formation of the supersaturated solid solution is insufficient, and even heating for 168 hours or more has no significant effect on the formation of the supersaturated solid solution. In the present invention, the melting point denotes a solidus temperature.

[0066] Additionally, rapid cooling from the solution annealing temperature is necessary to prevent excessive grain boundary reaction in the high-temperature range. In other words, quenching is required. This is because, although ductility improves due to grain boundary strengthening, if grain boundary reaction occurs in the high-temperature range, the hardness after aging treatment may be insufficient. Specifically, the cooling rate is preferably 10 °C / s or more, more preferably 50 °C / s or more, and further preferably 150 °C / s or more. Conversely, the cooling rate is preferably low to avoid quenching cracks, dimensional changes, deformations, and the like.Therefore, the above cooling rate, referred to as quenching, is unnecessary in a low-temperature range where no grain boundary reaction occurs and no spinodal segregation and / or excessive order formation takes place, from the perspective of increasing the hardness of the precious metal alloy. For example, quenching is not required if the cooling rate is in a temperature range below 250 °C. This suppresses or reduces the occurrence of quenching cracks and the like, and thus, for example, air cooling can be used in a temperature range of 200 °C or less after quenching to 200 °C. The endpoint of cooling in the solution annealing treatment is preferably room temperature. (B-3) Aging treatment

[0067] Spinodal demixing and order formation in the precious metal alloy of the present invention progress through an aging treatment of the supersaturated solid solution, formed by the solution annealing treatment, in a temperature range which is lower than the spinodal demixing temperature and the order-disorder transformation temperature.

[0068] The heating temperature, a condition of the aging treatment for the precious metal alloy in the supersaturated solid solution, is preferably 250 °C or more and 800 °C or less. A temperature below 250 °C hinders the progression of spinodal segregation or ordering. A temperature above 800 °C notably causes material softening due to grain boundary reactions. The heating temperature is more preferably 350 °C or more and 650 °C or less. The heating time during the aging treatment is preferably 0.001 hours or more and 168 hours or less. A heating time of less than 0.001 hours results in insufficient conversion and leads to hardness variation, while a treatment for 168 hours or more results in poor productivity and increased production costs. The cooling method after completion of the aging treatment is not particularly restricted.The precious metal alloy with high hardness of the present invention can be obtained through this aging treatment. (B-4) Other heat treatment step

[0069] The production of the precious metal alloy of the present invention comprises the aforementioned solution annealing and aging treatments as essential steps and may include other heat treatment steps. Examples of such other heat treatment steps include homogenization, two-phase treatment, and intermediate annealing. However, these other heat treatments have no effect on influencing the progress of spinodal segregation or ordering. Therefore, these heat treatments are optional.

[0070] A homogenization treatment is performed on a precious metal alloy produced by melt casting to create a metallic texture with a uniform distribution of element concentrations within the alloy. The homogenization treatment is a heat treatment in which a precious metal alloy is heated to a high temperature below its melting point for an extended period (preferably 0.1 hours or more and 72 hours or less).

[0071] A two-phase treatment is a treatment for forming a two-phase texture state, which provides the best workability in the alloy system of the present invention, and is a heat treatment carried out to facilitate warm forming and cold forming. The heating temperature in the two-phase treatment is 700 °C or more and 900 °C or less, and more preferably 750 °C or more and 850 °C or less. The heating time is 0.1 hours or more and 10 hours or less, and more preferably 0.2 hours or more and 2 hours or less.

[0072] Intermediate annealing is a heat treatment performed during the processing described below, such as warm forming or cold forming, in which stresses accumulate in a block or the like of a precious metal alloy. Intermediate annealing is a treatment to reduce the material's strength in order to restore its workability. The heating temperature during intermediate annealing is 700 °C or more and 900 °C or less, and more preferably 750 °C or more and 850 °C or less. The heating time is 0.1 hours or more and 10 hours or less, and more preferably 0.2 hours or more and 2 hours or less. (B-5) Processing step of the present precious metal alloy according to the invention

[0073] The precious metal alloy of the present invention can be processed into various shapes depending on its intended use by at least one processing step before the aging treatment is carried out. Examples of such processing steps include hot forming, semi-hot forming, cold forming, skin forming, straightening, coiling, and bending. Hot forming makes it possible to break up a solidified texture and eliminate defects such as cavities in a provided precious metal alloy block. Semi-hot forming and cold forming are significant not only for changing the overall shape of the alloy but also for controlling the shape of the crystal grains. If semi-hot forming or cold forming is carried out multiple times, the aforementioned intermediate annealing can be performed between the processing steps.

[0074] The precious metal alloy of the present invention, which exhibits improved ductility in its recrystallized texture, shows these effects particularly when subjected to semi-hot forming, cold forming, tempering, straightening, winding, or bending after solution annealing. After solution annealing (and prior to aging), the precious metal alloy is assumed to be processed into its final or near-final form for use. The precious metal alloy of the present invention exhibits improved ductility because the additive element B segregates and strengthens the grain boundaries, and therefore the alloy can be processed without defects even in a recrystallized texture state after solution annealing. It should be noted that the precious metal alloy of the present invention does not preclude processing after aging (after hardening).Even a precious metal alloy whose hardness is increased after aging treatment can be processed. Furthermore, after aging treatment, a precious metal alloy can also be subjected to final processing such as grinding / polishing, cutting, electrical discharge machining (EDM), pressing, bending, and straightening. Moreover, since the spinodal segregation and ordering that occur in the precious metal alloy of the present invention are reversible, the solution annealing and aging treatments can be repeated. Processing can also be carried out between combinations of multiple solution annealing and aging treatments. (B-6) Other method for producing the present precious metal alloy according to the invention

[0075] In the production process described above, the precious metal alloy of the present invention is obtained by subjecting an alloy ingot produced by a melting and casting process to solution treatment and aging treatment. In the present invention, while the solution annealing treatment is a primary heat treatment, in some cases a precious metal alloy can even be produced in a texture state corresponding to that of a supersaturated solid solution with a recrystallized texture without carrying out the solution treatment step. In such cases, the precious metal alloy of the present invention can be produced by subjecting this precious metal alloy to the aging treatment. The addition of B in the present invention can contribute to improving the ductility of the precious metal alloy in a recrystallized texture or a texture of a supersaturated solid solution prior to the aging treatment.

[0076] Examples of a step in obtaining a precious metal in a solid solution state without performing solution annealing include a liquid quenching process applied to a precious metal alloy in a molten state, and a step in applying a rapid heating and quenching process or a rapid solidification process using a laser or the like to a bulk metal. Alternatively, even various film formation processes such as welding, sputtering, plating, and thermal spraying can produce a precious metal alloy in a state that, in some cases, approximates a supersaturated solid solution state. The precious metal alloy produced by these processes can be aged without solution annealing to become the precious metal alloy of the present invention.These processes are useful for the production of near-net-shape precious metal alloys and layered / film-like precious metal alloys with coatings of high hardness or the like. Even if such a solution annealing treatment is not performed, the preferred conditions for aging treatment remain the same as described above. EXAMPLES

[0077] Specific examples of the present invention are described below. In these examples, a plurality of Au-Ni-Pd-Pt-based precious metal alloys were produced, with variations in the respective elemental composition of Au, Ni, Pd, Pt, B, Cu, and the metallic elements α and β. Subsequently, the ductility of the produced precious metal alloys was evaluated in a solution annealing state (in a recrystallized texture state), and the hardness was measured after an aging treatment. The production step and the evaluation procedures for the precious metal alloys of the examples were as follows. [Production of the precious metal alloy]

[0078] High-purity raw materials of the respective elements Au, Ni, Pd, Pt, B, Cu, and the metallic elements α and β were weighed and mixed to a predetermined composition. The resulting mixture was then placed in an alumina crucible, evacuated, and melted under a reduced-pressure argon atmosphere using high-frequency melting. The molten alloy was then poured into a copper mold (11 mm diameter, 70 mm length). The resulting ingot was then subjected to a homogenization treatment. This treatment was carried out taking into account the composition and melting point of the precious metal alloy, heating to a temperature of 800 °C to 1225 °C in a vacuum atmosphere (less than 5 × 10⁻⁵ m). -2 Pa). The heating time for the homogenization treatment was set to 16 hours in all cases.

[0079] The resulting ingot (11 mm diameter) from the homogenization treatment underwent a two-phase treatment (800 °C, 60 minutes, vacuum atmosphere), and the resulting piece was then cold-formed to a diameter of 6 mm by die forging, followed by cold forming using caliber rolls to a 2.5 mm square. A processing rate of 10 to 15% per pass was used for these processing steps, and intermediate annealing (800 °C, 60 minutes, vacuum atmosphere) was performed when the overall processing rate was within the range of 30 to 50%. In addition, wire drawing was carried out using a wire drawing machine to reduce the diameter to 0.6 mm, with a processing rate of 10 to 20% per pass. During this drawing process, intermediate annealing was performed within a range of 30 to 60% of the overall processing rate.Through these processing steps, the precious metal alloys of each of the examples and comparison examples were processed into wires with a diameter of 0.6 mm. These wires were then used as samples for evaluating ductility in a recrystallized state and for measuring hardness. [Solution annealing]

[0080] Next, the fabricated wire sample of each example was subjected to a solution annealing treatment. The solution annealing treatment was carried out in a vacuum atmosphere, with the heating temperature set to 925 °C to 1225 °C and the heating time set to 1 to 5 minutes. After heating, the sample was placed in a water cooling bath (20 ± 5 °C) within 3 seconds and cooled to room temperature. [Aging treatment]

[0081] After solution annealing, each wire sample underwent a bending test to assess ductility, followed by an aging treatment. The bending test procedure is described later. For the aging treatment, the heating temperature was set to 450 °C to 600 °C, and the sample was held at each aging temperature for 1 hour and then water-cooled to room temperature. Following aging, such a sample was embedded in a resin to remove an oxidized layer and residual stresses on a surface layer due to thermal expansion, and to obtain a hardness test specimen. The sample was then subjected to coarse polishing (#500, #800, #1200) and mirror polishing in 1 µm and 1 / 4 µm diamond suspensions.

[0082] In the present embodiment, Au-Ni-Pd-Pt-based precious metal alloys were produced as examples, and simultaneously, Au-Ni-Pd-Pt alloys, each with a composition similar to that of the corresponding example but without the addition of B, were produced as comparative examples. Additionally, Au-Ni-Pd-Pt-based alloys, each with a composition similar to that of the corresponding example but with the addition of C (carbon) instead of B, were also produced as comparative examples. Furthermore, Au-Ni-Pd-Pt-based precious metal alloys, each with the addition of Ag or Zr, were produced as comparative examples in connection with the grain boundary strengthening effect caused by grain boundary segregation. [Examination of material texture by X-ray diffraction analysis]

[0083] Furthermore, X-ray diffraction analysis was performed on each precious metal alloy after aging treatment to confirm the occurrence of spinodal segregation and the generation of an ordered phase. The conditions for the X-ray diffraction analysis of the respective samples were as follows: This X-ray diffraction analysis was performed on a precious metal alloy that had undergone solution annealing (solution-annealed material) and on a precious metal alloy that had undergone aging treatment (aged material) after solution annealing, so that the occurrence of spinodal segregation and order formation due to the aging treatment could be confirmed, and the analysis results compared.Each sample for X-ray diffraction analysis was prepared to improve the signal-to-noise ratio of an ordered reflection peak with low diffraction intensity by casting a precious metal alloy to a diameter of 30 mm × 100 mm and subjecting it to the same homogenization and two-phase treatment (800 °C, 60 minutes) as described above, as well as the same solution annealing (with a holding time of 10 minutes) and aging treatment as described above, after cold die forging to a diameter of 22 mm. X-ray diffraction analysis was then performed on the sample that had undergone the solution annealing treatment (solution-annealed material) and on the sample that had undergone the aging treatment (aged material). Resin embedding and polishing of each analytical sample was performed using the same procedure as described above. [General Terms and Conditions] - Sample size: φ22 mm × 2 mm - X-ray diffraction apparatus: SmartLab, manufactured by Rigaku - Target: Cu anode - Optical system detector: focusing optical system semiconductor detector (HyPix-3000) - Voltage-current: 40 kV-30 mA - Length-limiting gap: 10 mm (a) Confirmation of spinodal separation (lateral ligament peak) - 2θ scanning range: 20° to 130 - 2θ step size (°): 0.0012 - 2θ scan rate (° / min): 6 (b) Confirmation of ordered phase (ordered peak) 2θ scanning range: 20° to 38° 2θ step size (°): 0.0132 2θ scan rate (° / min): 1.3

[0084] In the investigation of spinodal separation by X-ray diffraction, it was determined whether one or more lateral ligament peaks appeared on one or both sides around both ends (by ± 0.5 to 3° with respect to 2θ) of a main peak observed with respect to each of the Miller indices {111}-plane, {200}-plane, {220}-plane, and {311}-plane in an X-ray diffraction profile obtained under the above conditions. An evaluation was performed as follows: A case in which one or more lateral ligament peaks appeared was considered to indicate the occurrence of spinodal separation, and a case in which no lateral ligament peaks appeared at all was considered to indicate the absence of spinodal separation.

[0085] In the investigation of an ordered phase using X-ray diffraction, the presence or absence of an ordered reflection peak generated around 2θ = 30° to 35° was determined. Specifically, a case in which the intensity of an ordered reflection peak was higher than the background of the corresponding region was considered the presence of an ordered phase, and a case in which the peak intensity was equal to or lower than the background was considered the absence of an ordered phase.

[0086] As an example of the results of the X-ray diffraction analysis carried out in the present embodiment, in Fig. 1 The X-ray diffraction patterns of the solution-annealed material and the aged material of the precious metal alloy (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) of Example 4-2 are shown. Fig. 1(a) presents an X-ray diffraction profile to confirm spinodal segregation, and Fig. 1(b) presents an X-ray diffraction profile to confirm an ordered phase. With reference to Fig. 1(a) In this precious metal alloy, clear peaks, identifiable as sideband peaks, were observed on both sides of a peak near 2θ = 40° to 42°, corresponding to the {111} plane, and a peak near 2θ = 46.5° to 48.5°, corresponding to the {200} plane. From this, it can be determined that spinodal demixing had occurred in this precious metal alloy. In addition, with reference to Fig. 1(b) In a region near 2θ = 31.5° to 34°, a peak (an ordered reflection peak) was observed which was definitely higher than the background. From this, it can be determined that order formation had also occurred in this precious metal alloy. [Examination of material texture by TEM / STEM analysis]

[0087] A TEM / STEM analysis was performed on the precious metal alloys produced in the present embodiment to confirm the occurrence of spinodal demixing and order formation. For the TEM / STEM analysis, samples (0.6 mm diameter) that had undergone aging treatment following solution annealing were first subjected to resin embedding polishing, and then samples for TEM / STEM analysis were prepared using a focused ion beam (FIB), on which the TEM / STEM analysis was then performed.

[0088] Fig. Figure 2 presents the results of the distribution measurement (Pt, Au, Pd: L-line, Ni: K-line) of each constituent element (Au, Ni, Pd, Pt) by STEM-EDS of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35). The electron incidence direction was parallel to the

[001] direction. Fig. 2. It was found that the material texture of the precious metal alloy of this embodiment exhibited a modulated texture with two regions, i.e., one in which the Au and Pd concentrations were relatively high, and the other in which the Au and Pd concentrations were relatively low, and that these regions were arranged alternately. The modulated texture does not exhibit a clear interface and was thus assumed to be due to spinodal demixing. Additionally, by comparison between the distribution pattern and the electron diffraction pattern, it was confirmed that the alloy has a modulated structure in the <010> -direction.

[0089] Fig. Figure 3 represents an electron beam diffraction pattern of the precious metal alloy of Example 4-2, which was obtained by TEM analysis (incident beam along the

[001] crystal zone axis). Fig. 3. It was confirmed that in this precious metal alloy not only a fundamental reflection with an fcc structure, but also a diffraction spot due to an ordered phase was observed. Based on the position of appearance, intensity, and plane spacing of the diffraction spot due to the ordered phase, it was assumed that the ordered phase has an L12 structure. In addition, since fringes in the <010> - and <100> -directions were observed at the 010 and 100 reciprocal grid points, indicating the presence of a modulated structure in the <010> -direction suggested. This result is considered to be consistent with the distribution measurement results obtained by STEM-EDS.

[0090] Similar results from the TEM / STEM analysis described above were also obtained for the other precious metal alloys. Therefore, it was confirmed that, in addition to X-ray diffraction analysis (diffraction methods), elemental analysis (composition mapping) and electron diffraction methods are also useful for determining spinodal demixing and order formation. [Confirmation of segregation in precious metal alloy after solution annealing]

[0091] Fig. Figure 4 presents the results of the STEM-EDS analysis, which was carried out near a high-angle grain boundary in a precious metal alloy of comparison example 2 (Au7.5-Ni41.25-Pd10-Pt41.25) after solution annealing. Fig. Figure 4 corresponds to the center of the photograph of a high-angle grain boundary. While the Pt and Ni concentrations near the grain boundary were low, the Au and Pd concentrations were high. This confirms that in the Au-Ni-Pd-Pt-based precious metal alloy of the present embodiment, Au and Pd tend to segregate near the grain boundary due to the solution annealing treatment.

[0092] Next, it was checked whether B segregated at the grain boundaries in each of the precious metal alloys of the samples to which B had been added. In the present embodiment, AES (Auger electron spectroscopy) analysis was performed to detect the light element B. A PHI-700 model manufactured by ULVAC-PHI was used as the AES apparatus. In the present embodiment, a precious metal alloy sample was impact fractured under ultra-high vacuum within the apparatus, and subsequently the fracture surface was observed, and AES analysis was performed on both an intergranular and a transgranular fracture surface. To prepare a sample for AES analysis, a precious metal alloy cast to a diameter of 11 mm was subjected to homogenization and two-phase treatment similar to those described above, and the result was cold-forged to a diameter of 6 mm.Subsequently, while intermediate annealing was performed, the result was cold-rolled to a 3.75 mm square potash thickness, followed by a solution anneal (1175 °C, 10 minutes), and the result was used as an evaluation specimen. The temperature and duration of this solution anneal were conditions intended to intentionally coarsen the crystal grains to facilitate the appearance of an intergranular fracture surface. Furthermore, an impact test was performed on the specimen in the cooled state (approximately -100 °C). As a result, both intergranular and transgranular fracture surfaces were obtained on the fracture surface after impact fracture, and the Auger spectra of both were measured.

[0093] As an example of the results of the AES analysis of the precious metal alloys of the examples, Fig. Five Auger spectra of the intergranular and transgranular fracture surfaces of the precious metal alloy of Example 4-2 (Au12.5-Ni37.325-Pd12.5-Pt37.325-B0.35) are presented. A comparison of the spectra of the intergranular and transgranular fracture surfaces revealed a peak near 185 eV only in the intergranular fracture surface. Since this peak near 185 eV was assumed to originate from B, it was assumed that B had segregated at the grain boundaries of the recrystallized texture. As a result, it was confirmed that the B added to the precious metal alloy tended to segregate at the grain boundaries. [Assessment of the ductility of the recrystallized texture]

[0094] The ductility of the wire after solution annealing was evaluated. The ductility evaluation test was performed on a wire sample with a diameter of 0.6 mm by conducting a bending test (bend radius 1 mm, bent at 180°) and measuring the bend angle until breakage. In the bending test, a measurement was performed five times on the same sample, and the average bend angle was evaluated as the bend value. For samples that did not break, even when bent at 180°, the bend value was considered to be 180°. [Measurement of hardness after aging treatment]

[0095] Hardness testing of a wire sample after aging treatment was performed at a test load of 0.075 kgf and at room temperature using a measuring instrument (HM-210, manufactured by Mitutoyo Corporation). The measurement results are shown in Table 1. A hardness measurement was taken at random at 15 points on each sample, and the average value was defined as the hardness value. The measurement positions in each sample were determined by selecting a plurality of crystal grains, and a measurement was taken in a non-grain boundary reaction region of each crystal grain, and as close to the center of the crystal grain as possible. In the hardness measurement results presented below, the aging treatment temperature at which the highest hardness was obtained among a plurality of aging temperatures set in the range of 450 °C to 600 °C is shown alongside the hardness value.However, the precious metal alloys in the examples showed a hardness of 500 Hv or more, which was the lowest.

[0096] Tables 1 to 4 show the compositions, ductility assessment results, and hardness measurement results for the precious metal alloys of the examples and comparison examples produced in the present embodiment. These tables also show the presence or absence of a modulated texture (sideband peaks) and an ordered phase (ordered peaks), as determined by the aforementioned X-ray diffraction analysis. In addition to the alloy compositions listed in Tables 1 to 4, the analyzed values ​​obtained from the wire samples of the respective examples are also shown. The alloy compositions were analyzed by ICP-OES analysis, and for samples where the B concentration was expected to be trace (0.05 atomic percent or less), a combination of ICP-MS analysis and ICP-MS was also used to determine the B concentration.In these tables, to confirm the effects of adding the essential additive element B and optional additive elements such as Cu and the metallic elements α and β, the precious metal alloys are categorized into examples and comparative examples, while being grouped according to precious metal alloys with similar concentrations of the essential metallic elements (Au, Ni, Pd, Pt).

[0097] The examination of the results shown in Tables 1 to 4 confirmed the prerequisite that all of the Au-Ni-Pd-Pt-based alloys in the respective examples and comparison examples, with the exception of those that could not undergo solution annealing due to cracks that occurred during processing, exhibited spinodal segregation and / or ordering, and confirmed a hardness of 500 Hv or more. It is therefore considered that the Au-Ni-Pd-Pt-based alloy on which the present applicant has focused has suitable potential as a high-hardness precious metal alloy.

[0098] Therefore, the effect of B, i.e., an essential additive element for improving the ductility of the recrystallized texture after solution annealing, is now investigated, which is the objective of the present invention. The tables showed that the Au-Ni-Pd-Pt-based alloys of each of the examples with added B, compared with the Au-Ni-Pd-Pt-based alloys of the comparison examples without added B, clearly exhibited an increase in the bending angle. For example, this point is considered clear with respect to Examples 1, 2, 3-1 to 3-5, and 4-1 to 4-4, which are shown in Table 1, and comparison examples corresponding to these examples. The same applies when comparing the combinations of examples and comparison examples shown in the other tables.Accordingly, it is considered to have been confirmed that the addition of B has the effect of improving the ductility of the recrystallized texture after solution annealing treatment.

[0099] However, the results in Table 1 confirm that the amount of B added to the Au-Ni-Pd-Pt-based alloy should also be within a specified range. If the amount of B added was too small (0.001 atomic percent), no effect on improving ductility was observed (Comparison Example 3). Furthermore, if B was added in excess, the ductility decreased in a two-phase texture state where the workability is initially good, leading to the formation of cracks during processing (Comparison Example 11-2).

[0100] Based on the results of the above-described investigation on the addition of B, it is assumed that in the Au-Ni-Pd-Pt-based alloy of the present invention, the added B segregates at the grain boundaries, which alters the grain boundary structure, and as a result increases the grain boundary strength to suppress grain boundary embrittlement, and thus improves the ductility of the recrystallized texture.

[0101] Regarding the effect of B, i.e., a key additive element of the Au-Ni-Pd-Pt-based alloy of the present invention, the results of investigations into the effect of additive elements used instead of B are shown as 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). Investigation of the evaluation results of these Au-Ni-Pd-Pt-based alloys in the comparative examples revealed that the addition of C and Zr did not improve the ductility of the recrystallized texture (comparative examples 4-2 and 4-4). Furthermore, when the concentration of added C was increased, grain boundary embrittlement became apparent, and even in the two-phase texture state, grain boundary cracks occurred, making processing difficult (comparative example 4-3).The addition of Ag, an element with a strong tendency to grain boundary segregation, was also attempted, but grain boundary embrittlement became visible, and even in the two-phase texture state grain boundary cracks occurred, making processing difficult (Comparison Example 3-5).

[0102] The effect of improving the ductility of the Au-Ni-Pd-Pt-based alloy in a recrystallized texture state by the addition of B was also confirmed by a tensile test at room temperature. Fig. Figure 6 presents the results of the room-temperature tensile test performed 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 test was performed on specimens (0.6 mm diameter × 50 mm length) after solution annealing under conditions of a chuck spacing of 25 mm and a crosshead speed of 10 mm / min. As in Fig. 6 confirmed that a significant improvement in elongation was observed when B was added, and thus it is assumed that the ductility was significantly improved.

[0103] Next, the effects of the optional additions of copper and metal elements α and β are investigated. Table 3 shows the results for the precious metal alloys obtained by adding copper to the Au-Ni-Pd-Pt-based alloy. Table 3 confirmed that spinodal segregation and order formation also occurred in the Au-Ni-Pd-Pt-based alloys with added copper. These precious metal alloys exhibited good hardness of 500 Hv or more after aging treatment.

[0104] Furthermore, the effect of adding boron (B) to the Au-Ni-Pd-Pt-based alloys with added copper (Cu) was also confirmed. It can be confirmed that the precious metal alloys without added B in the comparison examples exhibited poor ductility and grain boundary embrittlement, and that adding B, as in each of the examples, improved ductility and suppressed grain boundary embrittlement.

[0105] Table 4 shows the effects of adding the metal element α (In, Sn, Sb) and the metal element β (Ti, Zr, Hf, Al). First, it was confirmed that none of these additional elements inhibited the occurrence of spinodal segregation and order formation in the Au-Ni-Pd-Pt-based alloy. It was also observed that the addition of at least one of these elements increased hardness, independent of the addition of β. Depending on the composition of the precious metal alloy, this increase in hardness could, in some cases, be 100 Hv or more.

[0106] It was also confirmed that the addition of B in the Au-Ni-Pd-Pt-based alloys containing metallic elements α and β also showed an effect of improving ductility. Furthermore, a comparison between Example 24 and Example 26 in Table 3 shows that metallic elements α and β were also effective in increasing the hardness of the Au-Ni-Pd-Pt-based alloys containing the optional addition of copper. The metallic elements α and β can be added either individually or together, and multiple types of metallic elements can also be added. INDUSTRIAL APPLICABILITY

[0107] As previously described, the present invention relates to a precious metal alloy with increased hardness through spinodal demixing and / or ordering, in which the ductility is improved in a recrystallized texture formed by solution annealing or the like. In the present invention, it is considered that for grain boundary embrittlement caused by grain boundary segregation of Au and Pd after solution annealing, the addition of B causes B segregation at the grain boundaries, which modifies the grain boundary structure and, as a result, improves the grain boundary strength to suppress grain boundary embrittlement.

[0108] The precious metal alloy of the present invention is expected to be used in various applications requiring high hardness / high wear resistance, for example, in electrical / electronic materials such as styli and electrical contacts, medical tools, high-hardness coating elements, and the like. In these applications, processing after heat treatment, such as solution annealing, is considered an essential step. The precious metal alloy of the present invention exhibits good ductility after heat treatment, such as solution annealing, and can therefore be effectively provided for these applications.The precious metal alloy of the present invention can be produced by subjecting a precious metal alloy obtained by additive manufacturing techniques, metal powder injection molding techniques, rapid solidification techniques, laser heating technology, superimposed welding technology or coating techniques such as sputtering, thermal spraying and plating to an aging treatment, and is applicable for the above various uses. QUOTES INCLUDED IN THE DESCRIPTION

[0000] This list of documents cited by the applicant was automatically generated and is included solely for the reader's convenience. The list is not part of the German patent or utility model application. The DPMA accepts no liability for any errors or omissions. Cited patent literature

[0000] JP 2005-233967

[0009] JP 6997354

[0009] JP 2012-242184

[0009] JP 6372952

[0009] WO 2023 / 063156

[0009]

Claims

[1] Au-Ni-Pd-Pt-based precious metal alloy, comprising: 1.5 atomic percent or more and 47 atomic percent or less Au; 4 atomic % or more and 57 atomic % or less Ni; 1 atomic percent or more and 43.5 atomic percent or less Pd; and 6 atomic percent or more and 58.5 atomic percent or less Pt, wherein the Au-Ni-Pd-Pt-based precious metal alloy comprises 0.003 atomic % or more and 8 atomic % or less B. [2] Au-Ni-Pd-Pt-based precious metal alloy according to claim 1, further comprising more than 0.1 atomic % and 27.5 atomic % or less Cu. [3] 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, wherein the metal element α is at least one metal element selected from In, Sn and Sb. [4] 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, wherein the metal element β is at least one metal element selected from Al, Ti, Zr and Hf. [5] 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, wherein the metal element β is at least one metal element selected from Al, Ti, Zr and Hf. [6] Au-Ni-Pd-Pt-based precious metal alloy according to claim 1 or 2, wherein a material texture comprises a modulated texture by spinodal demixing. [7] Au-Ni-Pd-Pt-based precious metal alloy according to claim 1 or 2, wherein a material texture comprises an ordered phase. [8] Au-Ni-Pd-Pt-based precious metal alloy according to claim 6, wherein the material texture comprises an ordered phase. [9] 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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