EP3485050B1 - Copper-nickel-tin alloy, method for the production and use thereof - Google Patents

Description

The invention relates to a copper-nickel-tin alloy with excellent castability, hot workability and cold workability, high resistance to abrasive wear, adhesive wear and fretting wear and improved corrosion resistance and stress relaxation resistance according to the preamble of one of claims 1 or 2, a method their production according to the preamble of claims 9 or 10 and their use according to the preamble of claims 16 to 18.

Due to their good strength properties, their good corrosion resistance and conductivity for heat and electricity, the binary copper-tin alloys are of great importance in mechanical engineering and vehicle construction as well as in many areas of electronics and electrical engineering.

This group of materials has a high resistance to abrasive wear. In addition, the copper-tin alloys ensure good sliding properties and high fatigue strength, which is why they are extremely suitable for sliding elements in engine and vehicle construction as well as in general mechanical engineering.

The copper-nickel-tin alloys have compared to the binary copper-tin materials improved mechanical properties such as hardness, tensile strength and yield point. The increase in the mechanical characteristics is achieved through the hardenability of the Cu-Ni-Sn alloys.

In addition to the importance of the ratio of the elements nickel and tin for the temperature at which spontaneous spinodal demixing occurs in the Cu-Ni-Sn alloys, the precipitation processes are essential for the adjustment of the properties of this group of materials.

The presence of discontinuous precipitations, particularly at the grain boundaries of the structure of Cu-Ni-Sn alloys, is associated in the literature with a deterioration in toughness properties under dynamic stress.

So it is in the pamphlet DE 0 833 954 T1 proposed a Cu-Ni-Sn spinodal continuous casting alloy with 8 to 16 wt% Ni, 5 to 8 wt% Sn and optionally with up to 0.3 wt% Mn, up to 0.3 wt. -% B, up to 0.3 wt% Zr, up to 0.3 wt% Fe, up to 0.3 wt% Nb and up to 0.3 wt% Mg without wrought processing to manufacture. After solution heat treatment of the as-cast condition and after spinodal aging, the alloy must be rapidly cooled by water quenching in order to obtain a spinodal segregated structure without discontinuous precipitations.

In the pamphlet DE 23 50 389 C with reference to a Cu-Ni-Sn alloy with 2 to 98% by weight Ni and 2 to 20% by weight Sn, on the other hand, it is stated that cold forming with at least a degree of forming of ε=75% must be carried out in order to to be able to prevent the formation of brittle discontinuous precipitations during aging.

In the pamphlet DE 691 05 805 T2 are the difficulties that occur in the industrial large-scale production of semi-finished products and components made of copper-nickel-tin alloys. The occurrence of Sn-rich segregations, particularly at the grain boundaries of the cast structure, severely limits the possibility of further processing being economical. The Sn-rich segregations, which cannot be easily eliminated even by thermomechanical processing of the as-cast Cu-Ni-Sn alloys, prevent a homogeneous distribution of the alloying elements in the matrix. However, this is a mandatory prerequisite for the hardenability of this group of materials. It is therefore proposed to finely atomize the melt of a copper alloy containing 4 to 18% by weight of Ni and 3 to 13% by weight of Sn and to collect the sprayed particles on a collecting surface. Subsequent rapid cooling is intended to counteract the formation of the Sn-rich grain boundary segregations.

From the pamphlet DE 41 26 079 C2 It is known that a number of copper alloys cannot be produced with the conventional method of ingot casting with subsequent hot forming and cold forming with intermediate annealing, or can only be produced with poor economics because hot forming is difficult due to grain boundary precipitation, segregation or other inhomogeneities.

These copper alloys also include the copper-nickel-tin materials. Therefore, to ensure cold working of the as-cast condition of such alloys, a strip casting process with precise control of the solidification rate of the melt is recommended.

As a result of increasing operating temperatures and operating pressures in modern engines, machines, systems and aggregates, a wide variety of mechanisms of damage to the individual system elements occur. So there is more and more the need, especially in the material-side and constructive design of sliding elements and connectors, in addition to the Types of sliding wear also to consider the mechanism of oscillating wear damage.

Dynamic friction wear, also known as fretting in technical terms, is friction wear that occurs between oscillating contact surfaces. In addition to the geometric wear or volume wear of the components, the reaction with the surrounding medium leads to fretting corrosion. The damage to the material can significantly reduce the local strength in the wear zone, in particular the fatigue strength. Fatigue cracks can emanate from the damaged component surface, leading to fatigue fractures/friction fractures. In the case of fretting corrosion, the fatigue strength of a component can fall well below the fatigue limit value of the material.

The mechanism of oscillating friction wear is significantly different from the types of sliding wear with unidirectional movement. In particular, the effects of corrosion are particularly pronounced in the case of dynamic friction wear.

From the pamphlet DE 10 2012 105 089 A1 shows the damage consequences of the dynamic friction wear of plain bearings. To ensure a stable position of the plain bearings, they are pressed into the bearing mount. During the pressing process, a high level of stress is built up on the plain bearing, which is further increased by the increasing loads, thermal expansion and dynamic shaft loads in modern engines. As a result of the excessive stress, changes in the geometry of the plain bearing can occur, which reduce the original bearing overhang. This allows micro-movements of the plain bearing relative to the bearing mount. Due to these cyclical relative movements with a low oscillation range at the contact surfaces between the bearing and the bearing mount, Dynamic friction wear/fretting corrosion/fretting of the back of the plain bearing. The consequence is the initiation of cracks and ultimately the frictional fatigue fracture of the plain bearing.

The results of fretting tests with various plain bearing materials indicate that Cu-Ni-Sn alloys in particular with a Ni content of more than 2% by weight, as is the case with spinodally hardening copper-nickel-tin alloys, have an insufficient Possess resistance to fretting wear.

In motors and machines, electrical connectors are often placed in an environment in which they are subjected to mechanical vibrational movements. If the elements of a connection arrangement are located on different assemblies which, as a result of mechanical loads, carry out movements relative to one another, a corresponding relative movement of the connection elements can occur. These relative movements lead to oscillating wear and fretting corrosion of the contact zone of the connector. Microcracks form in this contact zone, which greatly reduces the fatigue strength of the connector material. A failure of the connector due to fatigue fracture can be the result. Furthermore, due to the fretting corrosion, there is an increase in the contact resistance.

A combination of the material properties of wear resistance, ductility and corrosion resistance is therefore decisive for sufficient resistance to dynamic friction wear/fretting corrosion/fretting.

In order to increase the wear resistance of copper-nickel-tin alloys, it is necessary to add suitable wear supports to these materials. These wear carriers in the form of hard particles are intended to protect against the consequences of abrasive and adhesive wear. As hard particles different precipitation forms come into consideration in the Cu-Ni-Sn alloys.

In the pamphlet U.S. 6,379,478 B1 teaches a copper alloy for connectors with 0.4 to 3.0% by weight Ni, 1 to 11% by weight Sn, 0.1 to 1% by weight Si and 0.01 to 0.06% by weight - %P revealed. The fine precipitations of the nickel silicides and nickel phosphides are said to ensure high strength and good stress relaxation resistance of the alloy.

For the production of a sliding layer on a base body made of steel is in the publication U.S. 2,129,197 A designates a copper alloy which is applied to the base body by build-up welding and contains 77 to 92% by weight Cu, 8 to 18% by weight Sn, 1 to 5% by weight Ni, 0.5 to 3% by weight Si and 0.25 to 1 wt% Fe. The silicides and phosphides of the alloying elements nickel and iron should serve as wear carriers.

From the pamphlet U.S. 3,392,017 A is a low-melting copper alloy containing up to 0.4% by weight Si, 1 to 10% by weight Ni, 0.02 to 0.5% by weight B, 0.1 to 1% by weight P and 4 to 25 wt% Sn is known. This alloy can be applied in the form of cast rods as a filler metal to suitable metallic substrate surfaces. The alloy has improved ductility over the prior art and is machinable. Besides hardfacing, this Cu-Sn-Ni-Si-PB alloy can be used for spray deposition. The addition of phosphorus, silicon and boron is intended to improve the self-fluxing properties of the molten alloy and the wetting of the substrate surface and make the use of an additional flux superfluous.

The teaching disclosed in this publication specifies a particularly high P content of 0.2 to 0.6% by weight with a mandatory Si content of the alloy of 0.05 up to 0.15% by weight. This underlines the superficial demand for the self-flowing properties of the material. With this high P content, the hot workability of the alloy will be poor and the spinodal segregability of the structure will be insufficient.

According to the pamphlet U.S. 4,818,307 A For example, the size of the hard particles precipitated in a copper-based alloy has a major impact on its wear resistance. Complex silicide formations/boride formations of the elements nickel and iron, which reach a size of 5 to 100 µm, increase the wear resistance of a copper alloy with 5 to 30% by weight Ni, 1 to 5% by weight Si, 0 .5 to 3 wt% B and 4 to 30 wt% Fe. The element tin is not contained in this material. This material is applied to a suitable substrate by build-up welding as a wear protection layer.

The pamphlet U.S. 5,004,581A describes the same copper alloy as the previous one U.S. 4,818,307 A with an additional content of tin in the content range of 5 to 15% by weight and/or of zinc in the content range of 3 to 30% by weight. The addition of Sn and/or zinc in particular increases the material's resistance to adhesive wear. This material is also applied to a suitable substrate as a wear-resistant layer by means of build-up welding.

However, the copper alloy according to the publications U.S. 4,818,307 A and U.S. 5,004,581A due to the required size of the silicide formations/boride formations of the elements nickel and iron of 5 to 100 μm only have a very limited cold formability.

The publication discloses a precipitation-hardenable copper-nickel-tin alloy U.S. 5,041,176 A out. This copper base alloy contains 0.1 to 10 wt% Ni, 0.1 to 10 wt% Sn, 0.05 to 5 wt% Si, 0.01 to 5 wt% Fe and 0.0001 to 1 % by weight of boron. This material contains dispersely distributed intermetallic phases of the Ni-Si system. The properties of the alloy are also explained using exemplary embodiments that have no Fe content.

In the pamphlet KR 10 2002 0 008 710 A (Abstract) it is stated that spinodal Cu-Ni-Sn alloys with a Sn content greater than 6% by weight are not hot workable. Sn-rich segregations at the grain boundaries of the cast structure of the Cu-Ni-Sn alloys are given as the reason. Therefore, for the disclosed Cu-Ni-Sn multicomponent alloy for high-strength wires and sheets, the composition becomes 1 to 8 wt% Ni, 2 to 6 wt% Sn, and 0.1 to 5 wt% of two or more Elements from the group Al, Si, Sr, Ti and B specified.

From the patent U.S. 5,028,282 A discloses a copper alloy with 6 to 25% by weight Ni, 4 to 9% by weight Sn and other additives with a content of 0.04 to 5% by weight (individually or together). These other additives are (in % by weight): 0.03 to 4% Zn, 0.01 to 0.2% Zr, 0.03 to 1.5% Mn, 0.03 to 0.7% Fe, 0.03 to 0.5% Mg, 0.01 to 0.5% P, 0.03 to 0.7% Ti, 0.001 to 0.1% B, 0.03 to 0.7% Cr, 0.01 to 0.5% Co.

It is stated that the alloying elements Zn, Mn, Mg, P and B are added to the melt of the alloy for deoxidation. The elements Ti, Cr, Zr, Fe and Co have a grain-refining and strength-enhancing function.

By alloying with metalloids such as boron, silicon and phosphorus, the reduction in processing technology, which is important high base melting temperature. This is why these alloying additions are used in particular in the field of wear-resistant coating materials and high-temperature materials, which include, for example, the alloys of the Ni-Si-B and Ni-Cr-Si-B systems. In these materials, the alloying elements boron and silicon in particular are responsible for the strong reduction in the melting temperature of hard nickel-based alloys, which is why their use as self-fluxing hard nickel-based alloys becomes possible.

In the exposition DE 20 33 744 B contains important information on another function of the alloying element boron in Si-containing metallic melts. Accordingly, the addition of boron causes the oxides formed in the melt to break down and boron silicates to form, which rise to the surface of the coating layers and thus prevent further access of oxygen. In this way, a smooth surface of the coating layer can be realized.

In the pamphlet DE 102 08 635 B4 describes the processes in a diffusion solder joint in which intermetallic phases are present. Diffusion soldering is used to connect parts with different thermal expansion coefficients. When this soldering point is subjected to thermomechanical stress or during the soldering process itself, high stresses occur at the interfaces, which can lead to cracks, particularly in the vicinity of the intermetallic phases. As a remedy, mixing the solder components with particles is suggested, which compensates for the different expansion coefficients of the joining partners. For example, particles of boron silicates or phosphorus silicates can minimize the thermomechanical stress in the soldered joint due to their advantageous thermal expansion coefficients. In addition, these particles prevent the cracks that have already been induced from propagating.

In the exposition DE 24 40 010 B the influence of the element boron is emphasized in particular on the electrical conductivity of a silicon cast alloy with 0.1 to 2.0% by weight of boron and 4 to 14% by weight of iron. In this Si-based alloy, a high-melting Si-B phase is precipitated, which is referred to as silicon boride.

The silicon borides, which are mostly present in the modifications SiB ₃ , SiB ₄ , SiB ₆ and/or SiB _n determined by the boron content, differ significantly from silicon in their properties. These silicon borides have a metallic character, which is why they are electrically conductive. They have an extraordinarily high temperature resistance and oxidation resistance. The modification of SiB ₆ , which is preferred for sintered products, is used, for example, in the manufacture and processing of ceramics because of its very high hardness and its high abrasive wear resistance.

The usual wear-resistant hard alloys for surface coating consist of a relatively ductile matrix of the metals iron, cobalt and nickel with embedded silicides and borides as hard particles ( Knotek, O.; Lugscheider, E.; Reimann, H.: A contribution to the assessment of wear-resistant nickel-boron-silicon hard alloys. Journal of Materials Engineering 8 (1977) 10, pp. 331-335 ). The wide use of hard alloys of the Ni-Cr-Si, Ni-Cr-B, Ni-B-Si and Ni-Cr-B-Si systems is based on the increase in wear resistance caused by these hard particles. In addition to the silicides Ni ₃ Si and Ni ₅ Si ₂ , the Ni-B-Si alloys also contain the borides Ni ₃ B and the Ni-Si borides/Ni silicoborides Ni ₆ Si ₂ B. A certain inertia is also reported of silicide formation in the presence of the element boron. Further investigations of the alloy system Ni-B-Si led to the detection of the high-melting Ni-Si borides Ni ₆ Si ₂ B and Ni _4.29 Si ₂ B _1.43 ( Lugscheider, E.; Reimann, H.; Knotek, O.: The ternary system nickel-boron-silicon. Monthly magazines for chemistry 106 (1975) 5, pp. 1155-1165 ). These refractory Ni-Si borides exist in a relatively large homogeneity range towards boron and silicon.

In frequent applications, the element zinc is added to copper-nickel-tin alloys to lower the metal price. Functionally, the alloying element zinc causes the stronger formation of Sn-rich or Ni-Sn-rich phases from the melt. In addition, zinc increases the formation of precipitates in the Cu-Ni-Sn spinodal alloys.

Furthermore, in numerous applications, a certain Pb content is added to the copper-nickel-tin alloys to improve the emergency running properties and for better machinability.

The invention is based on the object of providing a high-strength copper-nickel-tin alloy which has excellent hot workability over the entire range of the nickel content and tin content of 2 to 10% by weight in each case. It should be possible to use a starting material for hot forming that has been produced by means of conventional casting processes without the absolute necessity of carrying out spray compacting or thin-strip casting.

After casting, the copper-nickel-tin alloy should be free of gas pores and shrinkage pores as well as stress cracks and be characterized by a structure with a uniform distribution of the tin-enriched phase components. In addition, the structure of the copper-nickel-tin alloy should already contain intermetallic phases after casting. This is important so that the alloy already has high strength, high hardness and sufficient wear resistance in the cast state. Furthermore, the as-cast state should already be characterized by high corrosion resistance.

The as-cast state of the copper-nickel-tin alloy should not first have to be homogenized by means of a suitable annealing treatment in order to be able to produce sufficient hot workability.

With regard to the processing properties of the copper-nickel-tin alloy, there is the aim on the one hand that its cold workability should not deteriorate significantly in comparison to the conventional Cu-Ni-Sn alloys despite the content of intermetallic phases. On the other hand, the requirement for a minimum degree of deformation of the cold forming performed should be dropped for the alloy. According to the state of the art, this is regarded as a prerequisite in order to be able to ensure spinodal segregation of the structure of the Cu-Ni-Sn materials without the formation of discontinuous precipitations.

Another requirement regarding the further processing of Cu-Ni-Sn materials, which correspond to the state of the art, relates to the cooling rate after the materials have been aged. It is considered necessary, after spinodal aging, to quickly cool the materials by means of water quenching in order to obtain a spinodal segregated structure without discontinuous precipitations. However, since dangerous internal stresses can form as a result of this cooling method after aging, the invention is based on the further object of preventing the formation of discontinuous precipitations on the alloy side during the entire production process, including aging.

By means of further processing, which includes at least one annealing or at least one hot forming and/or cold forming in addition to at least one annealing, a fine-grained structure containing hard particles with high strength, high heat resistance, high hardness, high resistance to stress relaxation and corrosion, sufficient electrical conductivity and a high Degree of resistance to the mechanisms of Set sliding wear and oscillating friction wear.

The invention is represented by the features of one of claims 1 or 2 in relation to a copper-nickel-tin alloy, by the features of claims 9 or 10 in relation to a production method and by the features of claims 16 to 18 in relation to a use. The further dependent claims relate to advantageous configurations and developments of the invention.

• 2,0 bis 10,0 % Ni,
• 2,0 bis 10,0 % Sn,
• 0,01 bis 1,5 % Si,
• 0,01 bis 1,0 % Fe,
• 0,002 bis 0,45 % B,
• 0,001 bis 0,15 % P,
• wahlweise noch bis maximal 2,0 % Co,
• wahlweise noch bis maximal 2,0 % Zn,
• wahlweise noch bis maximal 0,25 % Pb,
• Rest Kupfer und unvermeidbare Verunreinigungen,
• dadurch gekennzeichnet,
  • dass das Verhältnis Si/B der Elementgehalte in Gew.-% der Elemente Silicium und Bor minimal 0,4 und maximal 8 beträgt;
  • dass nach dem Gießen in der Legierung folgende Gefügebestandteile vorliegen:
    • a) Eine Si-haltige und P-haltige metallische Grundmasse mit, bezogen auf das Gesamtgefüge,
    • a1) bis zu 30 Volumen-% ersten Phasenbestandteilen, die mit der Summenformel Cu_hNi_kSn_m ein Verhältnis (h+k)/m der Elementgehalte in Atom-% von 2 bis 6 aufweisen,
    • a2) bis zu 20 Volumen-% zweiten Phasenbestandteilen, die mit der Summenformel Cu_pNi_rSn_s ein Verhältnis (p+r)/s der Elementgehalte in Atom-% von 10 bis 15 aufweisen und
    • a3) einem Rest an Kupfer-Mischkristall;
    • b) Phasen, die, bezogen auf das Gesamtgefüge,
    • b1) mit 0,01 bis 10 Volumen-% als Si-haltige und B-haltige Phasen, welche als Siliziumboride und als Borsilikate und/oder als Borphosphorsilikate ausgebildet sind,
    • b2) mit 1 bis 15 Volumen-% als Ni-Si-Boride mit der Summenformel Ni_xSi₂B mit x = 4 bis 6,
    • b3) mit 1 bis 15 Volumen-% als Ni-Boride,
    • b4) mit 0,1 bis 5 Volumen-% als Fe-Boride,
    • b5) mit 1 bis 5 Volumen-% als Ni-Phosphide,
    • b6) mit 0,1 bis 5 Volumen-% als Fe-Phosphide,
    • b7) mit 1 bis 5 Volumen-% als Ni-Silizide,
    • b8) mit 0,1 bis 5 Volumen-% als Fe-Silizide und/oder Fe-reiche Teilchen im Gefüge enthalten sind, die einzeln und/oder als Anlagerungsverbindungen und/oder Mischverbindungen vorliegen und von Zinn und/oder den ersten Phasenbestandteilen und/oder den zweiten Phasenbestandteilen ummantelt sind;
  • dass beim Gießen die Si-haltigen und B-haltigen Phasen, welche als Siliziumboride ausgebildet sind, die Ni-Si-Boride, Ni-Boride, Fe-Boride, Ni-Phosphide, Fe-Phosphide, Ni-Silizide sowie die Fe-Silizide und/oder Fe-reichen Teilchen, die einzeln und/oder als Anlagerungsverbindungen und/oder Mischverbindungen vorliegen, Keime für eine gleichmäßige Kristallisation während der Erstarrung/Abkühlung der Schmelze darstellen, so dass die ersten Phasenbestandteile und/oder die zweiten Phasenbestandteile inselartig und/oder netzartig gleichmäßig im Gefüge verteilt sind;
  • dass die Si-haltigen und B-haltigen Phasen, welche als Borsilikate und/oder Borphosphorsilikate ausgebildet sind, zusammen mit den Phosphorsilikaten die Rolle eines verschleißschützenden und korrosionsschützenden Überzuges auf den Halbzeugen und Bauteilen der Legierung übernehmen.

The invention includes a high strength as-cast copper-nickel-tin alloy having excellent castability, hot workability and cold workability, high resistance to abrasive wear, adhesive wear and fretting wear, and improved corrosion resistance and stress relaxation resistance, consisting of (by weight %):

• 2.0 to 10.0% Ni,
• 2.0 to 10.0% Sn,
• 0.01 to 1.5% Si,
• 0.01 to 1.0% Fe,
• 0.002 to 0.45% B,
• 0.001 to 0.15% P,
• optionally up to a maximum of 2.0% Co,
• optionally up to a maximum of 2.0% Zn,
• optionally up to a maximum of 0.25% Pb,
• balance copper and unavoidable impurities,
• characterized,
  • that the Si/B ratio of the element contents in % by weight of the elements silicon and boron is a minimum of 0.4 and a maximum of 8;
  • that the following structural components are present in the alloy after casting:
    • a) A Si-containing and P-containing metallic matrix with, based on the overall structure
    • a1) up to 30% by volume of first phase components which, with the molecular formula Cu _h Ni _k Sn _m , have a ratio (h+k)/m of the element contents in atomic % of 2 to 6,
    • a2) up to 20% by volume of second phase components which, with the empirical formula Cu _p Ni _r Sn _s , have a ratio (p+r)/s of the element contents in atomic % of 10 to 15 and
    • a3) a remainder of copper mixed crystal;
    • b) phases which, in relation to the overall structure,
    • b1) with 0.01 to 10% by volume as Si-containing and B-containing phases, which are in the form of silicon borides and boron silicates and/or boron phosphorus silicates,
    • b2) with 1 to 15% by volume as Ni-Si borides with the molecular formula Ni _x Si ₂ B with x = 4 to 6,
    • b3) with 1 to 15% by volume as Ni borides,
    • b4) with 0.1 to 5% by volume as Fe borides,
    • b5) with 1 to 5% by volume as Ni phosphides,
    • b6) with 0.1 to 5% by volume as Fe phosphides,
    • b7) with 1 to 5% by volume as Ni silicides,
    • b8) the structure contains 0.1 to 5% by volume as Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds and of tin and/or the first phase components and/or or coated with the second phase constituents;
  • that during casting the Si-containing and B-containing phases, which are formed as silicon borides, the Ni-Si borides, Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and the Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds, represent nuclei for uniform crystallization during the solidification/cooling of the melt, so that the first phase components and/or the second phase components are island-like and/or are evenly distributed in the structure like a net;
  • that the Si-containing and B-containing phases, which are formed as boron silicates and/or boron phosphorus silicates, together with the phosphorus silicates, assume the role of a wear-resistant and anti-corrosive coating on the semi-finished products and components of the alloy.

Advantageously, the first phase components and/or the second phase components are contained with at least 1% by volume in the cast structure of the alloy.

Due to the uniform distribution of the first phase components and/or the second phase components in the form of islands and/or in the form of a network, the structure is free of segregations. Segregations of this type are understood to mean accumulations of the first phase components and/or the second phase components in the cast structure, which are formed as grain boundary segregations which, when the casting is subjected to thermal and/or mechanical stress, cause damage to the structure in the form of cracks that can lead to fracture. After casting, the structure is still free of gas pores, shrinkage pores, stress cracks and discontinuous precipitations of the (Cu, Ni)-Sn system.

In this variant, the alloy is in the as-cast state.

• 2,0 bis 10,0 % Ni,
• 2,0 bis 10,0 % Sn,
• 0,01 bis 1,5 % Si,
• 0,01 bis 1,0 % Fe,
• 0,002 bis 0,45 % B,
• 0,001 bis 0,15 % P,
• wahlweise noch bis maximal 2,0 % Co,
• wahlweise noch bis maximal 2,0 % Zn,
• wahlweise noch bis maximal 0,25 % Pb,
• Rest Kupfer und unvermeidbare Verunreinigungen,
• dadurch gekennzeichnet,
  • dass das Verhältnis Si/B der Elementgehalte in Gew.-% der Elemente Silicium und Bor minimal 0,4 und maximal 8 beträgt;
  • dass nach der Weiterverarbeitung der Legierung durch zumindest eine Glühung oder durch zumindest eine Warmumformung und/oder Kaltumformung nebst zumindest einer Glühung folgende Gefügebestandteile vorliegen:
    • A) Eine metallische Grundmasse mit, bezogen auf das Gesamtgefüge,
    • A1) bis zu 15 Volumen-% ersten Phasenbestandteilen, die mit der Summenformel Cu_hNi_kSn_m ein Verhältnis (h+k)/m der Elementgehalte in Atom-% von 2 bis 6 aufweisen,
    • A2) bis zu 10 Volumen-% zweiten Phasenbestandteilen, die mit der Summenformel Cu_pNi_rSn_s ein Verhältnis (p+r)/s der Elementgehalte in Atom-% von 10 bis 15 aufweisen und
    • A3) einem Rest an Kupfer-Mischkristall;
    • B) Phasen, die, bezogen auf das Gesamtgefüge,
    • B1) mit 2 bis 35 Volumen-% als Si-haltige und B-haltige Phasen, welche als Siliziumboride und als Borsilikate und/oder als Borphosphorsilikate ausgebildet sind, Ni-Si-Boride mit der Summenformel Ni_xSi₂B mit x = 4 bis 6, Ni-Boride, Fe-Boride, Ni-Phosphide, Fe-Phosphide, Ni-Silizide sowie als Fe-Silizide und/oder Fe-reichen Teilchen im Gefüge enthalten sind, die einzeln und/oder als Anlagerungsverbindungen und/oder Mischverbindungen vorliegen und von Ausscheidungen des Systems (Cu, Ni)-Sn ummantelt sind,
    • B2) mit bis zu 80 Volumen-% als kontinuierliche Ausscheidungen des Systems (Cu, Ni)-Sn im Gefüge enthalten sind,
    • B3) mit 2 bis 35 Volumen-% als Ni-Phosphide, Fe-Phosphide, Ni-Silizide sowie als Fe-Silizide und/oder Fe-reiche Teilchen im Gefüge enthalten sind, die einzeln und/oder als Anlagerungsverbindungen und/oder Mischverbindungen vorliegen, von Ausscheidungen des Systems (Cu, Ni)-Sn ummantelt sind und eine Größe von kleiner 3 µm aufweisen;
  • dass die Si-haltigen und B-haltigen Phasen, welche als Siliziumboride ausgebildet sind, die Ni-Si-Boride, Ni-Boride, Fe-Boride, Ni-Phosphide, Fe-Phosphide, Ni-Silizide sowie die Fe-Silizide und/oder Fe-reichen Teilchen, die einzeln und/oder als Anlagerungsverbindungen und/oder Mischverbindungen vorliegen, Keime für eine statische und dynamische Rekristallisation des Gefüges während der Weiterverarbeitung der Legierung darstellen, wodurch sich ein gleichmäßiges und feinkörniges Gefüges einstellt;
  • dass die Si-haltigen und B-haltigen Phasen, welche als Borsilikate und/oder Borphosphorsilikate ausgebildet sind, zusammen mit den Phosphorsilikaten die Rolle eines verschleißschützenden und korrosionsschützenden Überzuges auf den Halbzeugen und Bauteilen der Legierung übernehmen.

Furthermore, the invention includes a high-strength copper-nickel-tin alloy after annealing or after hot working and/or cold working plus annealing, with excellent castability, hot workability and cold workability, high resistance to abrasive wear, adhesive wear and fretting wear, and improved Corrosion resistance and stress relaxation resistance consisting of (in % by weight):

• 2.0 to 10.0% Ni,
• 2.0 to 10.0% Sn,
• 0.01 to 1.5% Si,
• 0.01 to 1.0% Fe,
• 0.002 to 0.45% B,
• 0.001 to 0.15% P,
• optionally up to a maximum of 2.0% Co,
• optionally up to a maximum of 2.0% Zn,
• optionally up to a maximum of 0.25% Pb,
• balance copper and unavoidable impurities,
• characterized,
  • that the Si/B ratio of the element contents in % by weight of the elements silicon and boron is a minimum of 0.4 and a maximum of 8;
  • that after further processing of the alloy by at least one annealing or by at least one hot forming and/or cold forming in addition to at least one annealing, the following structural components are present:
    • A) A metallic matrix with, based on the overall structure,
    • A1) up to 15% by volume of first phase components which, with the molecular formula Cu _h Ni _k Sn _m , have a ratio (h+k)/m of the element contents in atomic % of 2 to 6,
    • A2) up to 10% by volume of second phase components which, with the empirical formula Cu _p Ni _r Sn _s , have a ratio (p+r)/s of the element contents in atomic % of 10 to 15 and
    • A3) a remainder of copper mixed crystal;
    • B) phases which, in relation to the overall structure,
    • _B1 ) with ₂ to 35 vol to 6, Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles are contained in the structure, individually and/or as addition compounds and/or mixed compounds present and are covered by precipitations of the system (Cu, Ni)-Sn,
    • B2) up to 80% by volume as continuous precipitations of the (Cu, Ni)-Sn system are contained in the microstructure,
    • B3) are contained in the structure with 2 to 35% by volume as Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds , are encased by precipitations of the system (Cu, Ni)-Sn and have a size of less than 3 µm;
  • that the Si-containing and B-containing phases, which are formed as silicon borides, the Ni-Si borides, Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and the Fe silicides and/ or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds, represent nuclei for static and dynamic recrystallization of the structure during further processing of the alloy, resulting in a uniform and fine-grained structure;
  • that the Si-containing and B-containing phases, which are formed as boron silicates and/or boron phosphorus silicates, together with the phosphorus silicates, assume the role of a wear-resistant and anti-corrosive coating on the semi-finished products and components of the alloy.

Advantageously, the continuous precipitates of the (Cu,Ni)-Sn system are contained with at least 0.1% by volume in the structure of the further processed state of the alloy.

Even after further processing of the alloy, the structure is free of segregation. Such segregations are understood to mean accumulations of the first phase components and/or the second phase components in the structure, which are formed as grain boundary segregations, which cause damage to the structure in the form of cracks, particularly when the components are subjected to dynamic stress, which can lead to fracture.

After further processing, the structure of the alloy is free of gas pores, shrinkage pores and stress cracks. It should be emphasized as an essential feature of the invention that the microstructure in the further processed state is free of discontinuous precipitations of the (Cu,Ni)-Sn system. In this second variant, the alloy is in the processed state.

The invention is based on the consideration that a copper-nickel-tin alloy with Si-containing and B-containing phases and with phases of the systems Ni-Si-B, Ni-B, Fe-B, Ni-P, Fe-P, Ni-Si and is provided with other Fe-containing phases. These phases significantly improve the

Processing properties Castability, hot workability and cold workability. Furthermore, these phases improve the performance of the alloy by increasing strength and resistance to abrasive wear, adhesive wear and fretting wear. These phases additionally improve corrosion resistance and stress relaxation resistance as other performance properties of the invention.

The copper-nickel-tin alloy according to the invention can be produced by means of the sand casting process, shell mold casting process, investment casting process, full mold casting process, die casting process, lost foam process and permanent mold casting process or with the aid of continuous or semi-continuous casting method are produced.

Although the use of process-technically complex and cost-intensive primary shaping techniques is possible, it is not absolutely necessary for the production of the copper-nickel-tin alloy according to the invention. For example, the use of spray compacting or thin-strip casting can be dispensed with. The cast shapes of the copper-nickel-tin alloy according to the invention can be hot-formed directly over the entire range of Sn content and Ni content, for example by hot rolling, extrusion or forging, without carrying out a homogenization anneal that is absolutely necessary. Furthermore, it is noteworthy that after chill casting or continuous casting of the shapes from the alloy according to the invention, no complex forging processes or upsetting processes have to be carried out at elevated temperature in order to weld, ie close, pores and cracks in the material. This largely eliminates the processing restrictions that previously existed in the production of semi-finished products and components made of copper-nickel-tin alloys passed.

The metallic matrix of the structure of the copper-nickel-tin alloy according to the invention consists in the as-cast state with increasing Sn content of the alloy, depending on the casting process, from increasing proportions of tin-enriched phases, which are distributed evenly in the copper mixed crystal (a-phase). are.

These tin-enriched phases of the metallic matrix can be divided into first phase components and second phase components. The first phase components can be specified with the molecular formula Cu _h Ni _k Sn _m and have a ratio (h+k)/m of the element content in atom % of 2 to 6. The second phase components can be specified with the molecular formula Cu _p Ni _r Sn _s and have a ratio (p+r)/s of the element content in atom % of 10 to 15.

The alloy according to the invention is characterized by Si-containing and B-containing phases, which can be divided into two groups.

The first group relates to the Si-containing and B-containing phases, which are in the form of silicon borides and can be present in the modifications SiB ₃ , SiB ₄ , SiB ₆ and SiB _n . The "n" in the compound SiB _n indicates the high solubility of the element boron in the silicon lattice.

The second group of Si-containing and B-containing phases relates to the silicate compounds of borosilicates and/or borophosphorus silicates.

In the copper-nickel-tin alloy according to the invention, the structural proportion of the Si-containing and B-containing phases, which are in the form of silicon borides and borosilicates and/or borophosphorus silicates, is at least 0.01 and maximum 10% by volume.

The evenly distributed arrangement of the first phase components and/or second phase components in the structure of the alloy according to the invention results in particular from the effect of the Si-containing and B-containing phases, which are designed as silicon borides, and the Ni-Si borides with the empirical formula Ni _x Si ₂ B with x = 4 to 6, most of which are already precipitated in the melt. Subsequently, during the solidification/cooling of the melt, the Ni borides and Fe borides are precipitated, preferably on the silicon borides and Ni—Si borides already present. All of the boride compounds, which are present individually and/or as addition compounds and/or mixed compounds, serve as primary nuclei during the further solidification/cooling of the melt.

In the further course of the solidification/cooling of the melt, the Ni phosphides, Fe phosphides, Ni silicides, Fe silicides and/or the Fe-rich particles separate preferentially on the already existing primary nuclei of the silicon borides, Ni—Si borides and the Ni borides and Fe borides, which are present individually and/or as addition compounds and/or mixed compounds, as secondary nuclei.

The Ni-Si borides and the Ni borides are each contained in the structure at 1 to 15% by volume. The Ni phosphides and Ni silicides are each present with a structural proportion of 1 to 5% by volume. The Fe borides, Fe phosphides and the Fe silicides and/or Fe-rich particles each account for 0.1 to 5% by volume of the microstructure.

Thus, the structure contains the Si-containing and B-containing phases, which are formed as silicon borides, the Ni-Si borides with the empirical formula Ni _x Si ₂ B with x = 4 to 6 as well as the Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides, Fe silicides and/or Fe-rich particles individually and/or as addition compounds and/or mixed compounds. These phases are referred to below as crystallization nuclei.

Finally, the element tin and/or the first phase constituents and/or the second phase constituents of the metallic matrix preferentially crystallize in the regions of the nuclei, whereby the nuclei are coated by tin and/or the first phase constituents and/or the second phase constituents. These crystallization nuclei, which are encased by tin and/or the first phase components and/or the second phase components, are referred to below as hard particles of the first class.

In the cast state of the alloy according to the invention, the hard particles of the first class have a size of less than 80 μm. Advantageously, the size of the first-class hard particles is less than 50 μm.

As the Sn content of the alloy increases, the island-like arrangement of the first phase components and/or the second phase components changes into a network-like arrangement in the structure.

In the cast structure of the copper-nickel-tin alloy according to the invention, the first phase components can have a proportion of up to 30% by volume. The second phase components take on a structural proportion of up to 20% by volume. Advantageously, the first phase components and/or the second phase components are contained with at least 1% by volume in the as-cast structure of the alloy.

As a result of the addition of the alloying element boron, during the casting of the alloy according to the invention an inhibited and thus only incomplete formation of the phosphides and silicides. For this reason, some phosphorus and silicon remain dissolved in the as-cast metal matrix.

The conventional copper-nickel-tin alloys have a relatively large solidification interval. This large solidification interval increases the risk of gas absorption during casting and causes uneven, coarse, mostly dendritic crystallization of the melt. The result is often gas pores and coarse Sn-rich segregations, at the phase boundary of which shrinkage pores and stress cracks often occur. In this group of materials, the Sn-rich segregations also occur preferentially at the grain boundaries.

The combined content of boron, silicon and phosphorus activates various processes in the melt of the alloy according to the invention, which significantly change its solidification behavior compared to conventional copper-nickel-tin alloys.

The elements boron, silicon and phosphorus perform a deoxidizing function in the melt of the invention. By adding boron and silicon, it is possible to reduce the phosphorus content without reducing the intensity of the deoxidation of the melt. This measure makes it possible to suppress the disadvantageous effects of sufficient deoxidation of the melt by means of a phosphorus additive. A high P content would further extend the already very large solidification interval of the copper-nickel-tin alloy, which would increase the susceptibility to porosity and segregation of this type of material. The adverse effects of adding phosphorus are reduced by limiting the P content in the alloy of the present invention to the range of 0.001 to 0.15 wt%.

The lowering of the base melting temperature, particularly due to the element boron and the crystallization nuclei, lead to a reduction in the solidification interval of the alloy according to the invention. As a result, the cast state of the invention has a very uniform structure with a fine distribution of the individual phase components. Thus, segregations enriched with tin do not occur in the alloy according to the invention, particularly at the grain boundaries.

In the melt of the alloy according to the invention, the elements boron, silicon and phosphorus bring about a reduction in the metal oxides. The elements are oxidized themselves, usually rise to the surface of the castings and form a protective layer there as boron silicates and/or boron phosphorus silicates and phosphorus silicates, which protects the castings from gas absorption. Extraordinarily smooth surfaces of the castings made from the alloy according to the invention were found, which indicate the formation of such a protective layer. The structure of the as-cast state of the invention was also free of gas pores over the entire cross-section of the cast parts.

The advantages of introducing boron silicates and phosphorus silicates for avoiding stress cracks between phases with different thermal expansion coefficients during diffusion soldering were mentioned as part of the comments on the publications mentioned.

A basic idea of the invention is the transfer of the effect of boron silicates, boron phosphorus silicates and phosphorus silicates with regard to the adjustment of the different thermal expansion coefficients of the joining partners during diffusion soldering to the processes during casting, hot forming and thermal treatment of the copper-nickel-tin materials. Due to the wide solidification range of these alloys, it occurs between the staggered crystallizing Sn-poor and Sn-rich Structural areas to large mechanical stresses that can lead to cracks and pores. Furthermore, these damage characteristics can also occur during hot forming and high-temperature annealing of copper-nickel-tin alloys due to the different hot forming behavior and the different thermal expansion coefficients of the Sn-poor and Sn-rich microstructure components.

The combined addition of boron, silicon and phosphorus to the copper-nickel-tin alloy according to the invention causes, on the one hand, a structure with a uniform island-like and/or network-like distribution of the first phase components and/or the second phase components during the solidification of the melt by means of the effect of the crystallization nuclei Phase components of the metallic matrix. In addition to the crystallization nuclei, the Si-containing and B-containing phases formed during solidification of the melt, which are in the form of boron silicates and/or boron phosphorus silicates, ensure, together with the phosphorus silicates, the necessary adjustment of the thermal expansion coefficients of the first phase components and/or the second Phase components and the copper mixed crystal of the metallic matrix. This prevents the formation of pores and stress cracks between the phases with different Sn contents.

The alloy content of the copper-nickel-tin alloy according to the invention also causes a significant change in the grain structure in the cast state. It could be determined that a substructure with a grain size of the subgrains of less than 30 µm is formed in the primary cast structure.

Alternatively, the alloy according to the invention can be subjected to further processing by annealing or by hot working and/or cold working in addition to at least one annealing.

One possibility for further processing of the copper-nickel-tin alloy according to the invention consists in converting the castings into the final shape with the required properties by means of at least one cold forming together with at least one annealing.

Due to the uniform cast structure and the first-class hard particles precipitated therein, the alloy according to the invention already has high strength in the cast state. As a result, the castings have lower cold formability, which makes further processing more difficult. For this reason, it has proven to be advantageous to carry out a homogenization annealing of the castings before cold forming.

To ensure the ability of the invention to age, accelerated cooling after the homogenization annealing processes has proven to be advantageous. It has been shown that due to the inertia of the precipitation mechanisms and segregation mechanisms, cooling methods with a lower cooling rate can be used in addition to water quenching. The use of accelerated air cooling has thus proven to be just as practicable in order to sufficiently reduce the effect of the precipitation processes and demixing processes in the structure, which increase hardness and increase strength, during the homogenization annealing of the invention.

The outstanding effect of the crystallization nuclei for the recrystallization of the structure of the invention can be seen in the structure, which can be adjusted after cold forming by means of annealing in the temperature range from 170 to 880° C. and an annealing time of between 10 minutes and 6 hours. The extraordinarily fine structure of the recrystallized alloy enables further cold forming steps with a degree of forming ε of mostly over 70%. To this In this way, high-strength states of the alloy can be produced.

Due to these high degrees of cold forming that have become possible during the further processing of the invention, particularly high values for the tensile strength R _m , the yield point R _p0.2 and the hardness can be set. In particular, the level of the parameter R _p0.2 is significant for the sliding elements and guide elements. Furthermore, a high value of R _{p0.2 is} a prerequisite for the necessary spring properties of connectors in electronics and electrical engineering.

In the statements of numerous publications, which describe the state of the art with regard to the processing and the properties of the copper-nickel-tin materials, reference is made to the necessity of maintaining a minimum degree of cold deformation of, for example, 75% in order to avoid discontinuous precipitations of the (Cu,Ni)-Sn system in the structure.

In contrast, the structure of the alloy according to the invention remains free of discontinuous precipitations of the (Cu,Ni)-Sn system, regardless of the degree of cold working. It was thus possible to establish for particularly advantageous embodiments of the invention that the microstructure of the invention remains free of discontinuous precipitations of the (Cu,Ni)—Sn system even with extremely small degrees of cold deformation of less than 20%.

According to the state of the art, the conventional, spinodally separable Cu-Ni-Sn materials are considered to be very difficult or impossible to hot-work.

The effect of the crystallization nuclei could also be observed during the process of hot working the copper-nickel-tin alloy according to the invention. Mainly the crystallization nuclei are for it responsible for the fact that the dynamic recrystallization takes place favorably during the hot forming of the alloy according to the invention in the temperature range from 600 to 880°C. This results in a further increase in the uniformity and fine-grained structure.

Advantageously, the semi-finished products and components can be cooled after hot forming in calmed or accelerated air or with water.

As after the casting, an extraordinarily smooth surface of the parts could also be determined after the hot forming of the castings. This observation points to the formation of Si-containing and B-containing phases, which are in the form of boron silicates and/or boron phosphorus silicates, and of phosphorus silicates, which takes place in the material during hot forming. The silicates, together with the crystallization nuclei, also cause the different thermal expansion coefficients of the phases of the metallic matrix of the invention to be equalized during hot working. The surface of the hot-formed parts and the structure were free of cracks and pores, just as they were after casting.

Advantageously, at least one annealing treatment of the cast state and/or the hot-formed state of the invention can be carried out in the temperature range from 170 to 880°C for a period of 10 minutes to 6 hours, alternatively with cooling in calmed or accelerated air or with water.

One aspect of the invention relates to an advantageous method for further processing the as-cast condition or the hot-worked condition or the annealed cast condition or the annealed hot-worked condition State, which includes the implementation of at least one cold forming.

At least one annealing treatment of the cold-formed state of the invention can preferably be carried out in the temperature range from 170 to 880° C. for a period of 10 minutes to 6 hours, alternatively with cooling in calmed or accelerated air or with water.

Stress relief annealing/ageing annealing can advantageously be carried out in the temperature range from 170 to 550° C. with a duration of 0.5 to 8 hours.

After further processing of the alloy by at least one annealing or by at least one hot forming and/or cold forming together with at least one annealing, precipitates of the (Cu, Ni)-Sn system form preferably in the areas of the crystallization nuclei, as a result of which the crystallization nuclei are encased by these precipitations.

These crystallization nuclei, encased by precipitations of the (Cu,Ni)-Sn system, are referred to below as second-class hard particles.

As a result of the further processing of the alloy according to the invention, the size of the second-class hard particles decreases in comparison to the size of the first-class hard particles. In particular, with an increasing degree of cold forming, there is progressive comminution of the hard particles of the second class, since these, as the hardest components of the alloy, cannot support the change in shape of the surrounding metallic matrix. The resulting second-class hard particles and/or the resulting segments of the second-class hard particles have a size of less than 40 μm to even less than 5 μm, depending on the degree of cold forming.

The Ni content and the Sn content of the invention range between 2.0 and 10.0% by weight. A Ni content and/or a Sn content of less than 2.0% by weight would result in insufficient strength and hardness values. In addition, the running properties of the alloy would be insufficient under a sliding load. The alloy's resistance to abrasive and adhesive wear would not meet the requirements. With a Ni content and/or a Sn content of more than 10.0% by weight, the toughness properties of the alloy according to the invention would deteriorate rapidly, as a result of which the dynamic load capacity of the components made from the material is reduced.

The content of nickel and tin in the range of 3.0 to 9.0 wt. In this respect, the range of 4.0 to 8.0% by weight each for the content of the elements nickel and tin is particularly preferred for the invention.

It is known from the prior art with regard to Ni-containing and Sn-containing copper materials that the degree of spinodal segregation of the structure increases with an increasing Ni/Sn ratio of the element contents in % by weight of the elements nickel and tin. This is valid for a Ni content and a Sn content from about 2% by weight. As the Ni/Sn ratio decreases, the mechanism of precipitation formation in the (Cu,Ni)-Sn system becomes more important, which leads to a reduction in the proportion of the spinodally segregated microstructure. One consequence is, in particular, a more pronounced formation of discontinuous precipitations of the (Cu,Ni)-Sn system as the Ni/Sn ratio decreases.

One of the essential features of the copper-nickel-tin alloy according to the invention is the decisive suppression of the influence of the Ni/Sn ratio on the formation of discontinuous precipitations in the structure. Thus, it has been found that in the microstructure of the invention, largely independent of the Ni/Sn ratio and independent of the aging conditions, there is no precipitation of discontinuous precipitates of the (Cu,Ni)-Sn system.

In contrast, during further processing of the alloy according to the invention, up to 80% by volume of continuous precipitations of the (Cu,Ni)—Sn system are formed. Advantageously, the continuous precipitates of the (Cu,Ni)-Sn system are contained with at least 0.1% by volume in the structure of the further processed state of the alloy.

The element iron is added to the alloy according to the invention at 0.01 to 1.0% by weight. Iron contributes to increasing the proportion of crystallization nuclei and thus supports the fine-grain formation of the structure during the casting process. The Fe-containing hard particles in the structure increase the strength, hardness and wear resistance of the alloy. If the Fe content is below 0.01% by weight, these effects on the structure and the properties of the alloy are only observed to an insufficient extent. If the Fe content exceeds 1.0% by weight, the structure increasingly contains cluster-like accumulations of Fe-rich particles. The Fe content of these clusters would only be available to a lesser extent for the formation of crystallization nuclei and hard particles. In addition, the toughness properties of the invention would deteriorate. An Fe content of 0.02 to 0.6% by weight is advantageous. An iron content in the range from 0.06 to 0.4% by weight is preferred.

Due to the similarity relationship between the elements nickel and iron, Fe-Si borides and/or Ni-Fe-Si borides can also form in the structure of the alloy according to the invention in addition to the Ni-Si borides. The Ni-Fe-Si borides can be given with the empirical formula (Ni, Fe) _x Si ₂ B with x = 4 to 6.

In addition to the Fe borides and Fe phosphides, the structure of the invention also contains other Fe-containing phases.

As a result of the determined inertia in the precipitation of the Fe silicides and the dependence of the precipitation of the Fe silicides on the process conditions during the production and further processing of the alloy according to the invention, these other Fe-containing phases are present as Fe silicides and/or as Fe-rich particles structure.

The effect of the crystallization nuclei during the solidification/cooling of the melt, the effect of the crystallization nuclei as recrystallization nuclei and the effect of the silicate-based phases for the purpose of wear protection and corrosion protection can only reach a technically significant level in the alloy according to the invention if the silicon content is at least 0 .01% by weight and the boron content is at least 0.002% by weight. On the other hand, if the Si content exceeds 1.5% by weight and/or the B content exceeds 0.45% by weight, the casting performance deteriorates. The excessively high content of crystallization nuclei would make the melt significantly more viscous. In addition, reduced toughness properties of the alloy according to the invention would result.

The range for the Si content within the limits of 0.05 to 0.9% by weight is rated as advantageous. A silicon content of 0.1 to 0.6% by weight has proven to be particularly advantageous.

A content of 0.01 to 0.4% by weight is considered advantageous for the element boron. A boron content of 0.02 to 0.3 has proven particularly advantageous % by weight.

A lower limit for the element ratio of the elements silicon and boron has proven to be important in order to ensure a sufficient content of Ni-Si borides and of phases containing Si and B, which are in the form of borosilicates and/or borophosphorus silicates. For this reason, the minimum Si/B ratio of the element contents of the elements silicon and boron in % by weight of the alloy according to the invention is 0.4. The minimum Si/B ratio of the element contents of the elements silicon and boron in % by weight of 0.8 is advantageous for the alloy according to the invention. The minimum Si/B ratio of the element contents of the elements silicon and boron in % by weight of 1 is preferred.

Important for another important feature of the invention is the determination of an upper limit of 8 for the Si/B ratio of the element contents of the elements silicon and boron in % by weight. After casting, parts of the silicon are dissolved in the metallic matrix and bound in the hard particles of the first class.

During further thermal or thermomechanical processing of the cast state, at least a partial dissolution of the silicide components of the first-class hard particles occurs. This increases the Si content of the metallic matrix. If this exceeds an upper limit value, an excessive proportion, particularly of Ni silicides, is precipitated as the size increases. These would significantly reduce the cold formability of the invention.

For this reason, the maximum Si/B ratio of the element contents of the elements silicon and boron in % by weight of the alloy according to the invention is 8 or thermomechanical further processing of the cast state of the alloy-forming silicides to below 3 μm. Furthermore, this limits the content of silicides. Limiting the Si/B ratio of the element contents of the elements silicon and boron in % by weight to the maximum value of 6 has proven to be particularly advantageous in this respect.

The precipitation of the crystallization nuclei influences the viscosity of the melt of the alloy according to the invention. This circumstance underlines why the addition of phosphorus must not be dispensed with. Phosphorus causes the melt to be sufficiently thin despite the crystallization nuclei, which is of great importance for the castability of the invention. The phosphorus content of the alloy according to the invention is 0.001 to 0.15% by weight.

Below 0.001% by weight, the P content no longer contributes to ensuring sufficient castability of the invention. If the phosphorus content of the alloy assumes values above 0.15% by weight, then on the one hand an excessive Ni content is bound in the form of phosphides, which reduces the spinodal segregability of the structure. On the other hand, with a P content above 0.15% by weight, the hot workability of the invention would deteriorate significantly. For this reason, a P content of 0.01 to 0.15% by weight has proven particularly advantageous. A P content in the range from 0.02 to 0.09% by weight is preferred.

The alloying element phosphorus is very important for another reason. Together with the required maximum Si/B ratio of the element contents of the elements silicon and boron in % by weight of 8, it is due to the phosphorus content of the alloy that after further processing of the invention Ni phosphides, Silicides and Fe-silicides and / or Fe-rich particles individually and / or as Addition compounds and/or mixed compounds are present and are encased by precipitations of the system (Cu, Ni)-Sn, with a maximum size of 3 µm and with a content of 2 to 35% by volume in the structure.

These Ni phosphides, Fe phosphides, Ni silicides, Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds, are encased by precipitations of the (Cu,Ni)-Sn system and a maximum size of 3 µm are referred to below as hard particles of the third class.

In the structure of the further processed state of the particularly preferred embodiment of the invention, the hard particles of the third class even have a size of less than 1 μm.

On the one hand, these third-class hard particles supplement the second-class hard particles in their function as wear carriers. They increase the strength and hardness of the metallic matrix and thus improve the resistance of the alloy to abrasive wear. On the other hand, the third-class hard particles increase the alloy's resistance to adhesive wear. Finally, these third-class hard particles bring about a significant increase in the high-temperature strength and the stress relaxation resistance of the alloy according to the invention. This represents an important prerequisite for the use of the alloy according to the invention, in particular for sliding elements and components and connecting elements in electronics/electrical engineering.

Due to the content of hard particles of the first class in the structure of the cast state and of hard particles of the second and third class in the structure of the further processed state, the alloy according to the invention has the Character of a precipitation-hardenable material. Advantageously, the invention corresponds to a precipitation hardenable and spinodally demixable copper-nickel-tin alloy.

The sum of the element contents of the elements silicon, boron and phosphorus is advantageously at least 0.2% by weight.

The cast variant and the further processed variant of the alloy according to the invention can contain the following optional elements:
The element cobalt can be added to the copper-nickel-tin alloy according to the invention in a content of up to 2.0% by weight. Due to the similarity relationship between the elements nickel, iron and cobalt and due to the Si-boride-forming, boride-forming, silicide-forming and phosphide-forming properties of cobalt with regard to nickel and iron, the alloying element cobalt can be added in order to participate in the formation of the crystallization nuclei and the hard particles first , second and third class of the alloy. Thereby, the Ni content bound in the hard particles can be reduced. In this way it can be achieved that the Ni content that is effectively available in the metallic matrix for the spinodal segregation of the structure increases. With the addition of advantageously 0.1 to 2.0% by weight of Co, it is thus possible to increase the strength and hardness of the invention considerably.

The element zinc can be added to the copper-nickel-tin alloy according to the invention in a content of 0.1 to 2.0% by weight. It has been found that the alloying element zinc, depending on the Ni content and Sn content of the alloy, increases the proportion of the first phase components and/or second phase components in the metallic matrix of the invention, thereby increasing strength and hardness. They are to be held responsible for this Interactions between the Ni content and the Zn content. As a result of these interactions between the Ni content and the Zn content, a decrease in the size of the first and second class hard particles was also determined, which were thus formed in a more finely distributed manner in the structure. Below 0.1% by weight of Zn, these effects on the structure and mechanical properties of the invention could not be observed. At Zn levels above 2.0 wt%, the toughness properties of the alloy were lowered to a lower level. In addition, the corrosion resistance of the copper-nickel-tin alloy according to the invention deteriorated. Advantageously, a zinc content ranging from 0.1 to 1.5% by weight can be added to the invention.

Optionally, the copper-nickel-tin alloy according to the invention can have small amounts of lead, which are above the impurity limit, up to a maximum of 0.25% by weight. In a particularly preferred advantageous embodiment of the invention, the copper-nickel-tin alloy is free of lead except for any unavoidable impurities, which means that current environmental standards are met. In this context, lead contents of up to a maximum of 0.1% by weight of Pb are contemplated.

The formation of Si-containing and B-containing phases, which are in the form of boron silicates and/or boron phosphorus silicates, as well as phosphorus silicates, not only leads to a significant reduction in the content of pores and cracks in the structure of the alloy according to the invention. These silicate-based phases also assume the role of an anti-wear and anti-corrosion coating on the components.

During the adhesive wear stress of a component made of the copper-nickel-tin alloy according to the invention, the alloying element tin contributes to the formation of a so-called tribolayer between the sliding partners. Especially under mixed friction conditions this mechanism is significant when the emergency running properties of a material are increasingly important. The tribological layer leads to a reduction in the purely metallic contact surface between the sliding partners, which prevents the elements from welding or seizing.

Due to the increase in efficiency of modern engines, machines and aggregates, higher and higher operating pressures and operating temperatures are occurring. This can be observed in particular in the newly developed internal combustion engines, in which work is being done towards more and more complete combustion of the fuel. In addition to the increased temperatures in the combustion engine room, there is also the heat that occurs during the operation of the plain bearing systems. As a result of the high temperatures in storage operation, the parts made of the alloy according to the invention form, similar to casting and hot forming, phases containing Si and B, which are in the form of boron silicates and/or boron phosphorus silicates, as well as phosphorus silicates . These compounds further reinforce the tribolayer, which forms primarily due to the alloying element tin, resulting in increased adhesive wear resistance of the sliding elements made from the alloy according to the invention.

The alloy according to the invention thus ensures a combination of the properties of wear resistance and corrosion resistance. This combination of properties leads to a high resistance to the mechanisms of sliding wear, as required, and to a high material resistance to fretting corrosion. In this way, the invention is excellently suited for use as a sliding element and plug connector, since it has a high degree of resistance to sliding wear and oscillating wear, so-called fretting.

In addition to the important contribution made by the hard particles of the third class to increasing the resistance of the invention to the abrasive and adhesive mechanism of sliding wear, the hard particles of the third class make a significant contribution to increasing the fatigue strength. The hard particles of the third class, together with the hard particles of the second class, represent obstacles to the propagation of fatigue cracks, which can be introduced into the stressed component, particularly in the case of dynamic friction wear, the so-called fretting. The hard particles of the second and third class supplement in particular the anti-wear and anti-corrosion effect of the Si-containing and B-containing phases, which are designed as boron silicates and/or boron phosphorus silicates, and the phosphorus silicates with regard to increasing the resistance of the alloy according to the invention to dynamic friction wear, the so-called fretting.

High temperature strength and stress relaxation resistance are other important properties of an alloy used in higher temperature applications. A high density of fine precipitates is considered to be advantageous in order to ensure sufficiently high heat resistance and stress relaxation resistance. In the alloy according to the invention, such precipitations are the hard particles of the third class and the continuous precipitations of the (Cu,Ni)—Sn system.

Due to the uniform and fine-grained structure with largely no pores, no cracks and no segregation and the content of first-class hard particles, the alloy according to the invention has a high degree of strength, hardness, ductility, complex wear resistance and corrosion resistance even in the cast state. This combination of properties means that sliding elements and guide elements can be produced from the cast format. The cast state of the invention can also be used for the production of fitting housings and housings of water pumps, oil pumps and fuel pumps. In addition, the alloy of the present invention can be used for propellers, blades, propellers and hubs for shipbuilding.

The further processed variant of the invention can be used for areas of application with a particularly strong, complex and/or dynamic component stress.

Due to the outstanding strength properties and wear resistance as well as corrosion resistance of the copper-nickel-tin alloy according to the invention, another possible application comes into consideration. Thus, the invention is suitable for the metal objects in constructions for the rearing of organisms living in sea water (aquaculture). Furthermore, pipes, seals and connecting bolts which are required in the maritime and chemical industries can be produced from the invention.

The material is of great importance for the use of the alloy according to the invention for the production of percussion instruments. In particular, cymbals of high quality have hitherto been made from mostly tin-containing copper alloys by means of hot forming and at least one annealing before they are brought into their final shape, usually by means of a bell or bowl. The basins are then annealed again before their final machining takes place. The production of the different variants of the cymbal (eg ride cymbal, hi-hat, crash cymbal, China cymbal, splash cymbal and effect cymbal) therefore requires a particularly advantageous hot workability of the material, which is ensured by the alloy according to the invention. Within the range limits of the chemical composition of the invention, different structural proportions of the Phases of the metallic matrix and the different hard particles can be set in a very wide range. In this way it is already possible to influence the sound of the cymbals on the alloy side.

The invention can be used in particular for the production of composite plain bearings in order to be applied to a composite partner by means of a joining method. It is thus possible to produce a composite between disks, plates or strips of the invention and steel cylinders or steel strips, preferably made of heat-treated steel, by means of forging, soldering or welding with the optional implementation of at least one annealing in the temperature range from 170 to 880°C. Likewise, for example, composite bearing shells or composite bearing bushes can be produced by roll-bonding, inductive or conductive roll-bonding or by laser roll-bonding, also with the optional implementation of at least one annealing in the temperature range from 170 to 880°C.

As a result of the formation of the microstructure in the alloy according to the invention, there are further possibilities for the production of composite sliding elements such as composite bearing shells or composite bearing bushes. It is thus possible to apply a coating of tin or a Sn-rich material to a base body of the invention by means of hot-dip tinning or galvanic tinning, sputtering or the PVD process or CVD process, which serves as a running layer in storage operations.

In this way, high-performance composite sliding elements such as composite bearing shells or composite bearing bushes can also be produced as a three-layer system, with a bearing back made of steel, the actual bearing made of the alloy according to the invention and the overlay made of tin or the Sn-rich coating. This multi-layer system has a particularly beneficial effect on the adaptability and shrinkability of the plain bearing and improves the embedding ability of foreign particles and abrasive particles, whereby thermal or thermomechanical stress on the plain bearing does not result in damage due to the layer composite system being broken as a result of pore formation and cracking in the border area of the individual layers.

The great potential of copper-nickel-tin materials, particularly with regard to strength, spring properties and resistance to stress relaxation, can also be used for the field of tin-plated components, line elements, guide elements and connecting elements in electronics and electrical engineering by using the alloy according to the invention. The structure of the invention reduces the damage mechanism of pore formation and cracking in the boundary area between the alloy according to the invention and the tinning, even at elevated temperatures, which counteracts an increase in the electrical contact resistance of the components or even detachment of the tinning.

Machining of the semi-finished products and components made of the conventional wrought copper-nickel-tin alloys with a Ni content and Sn content of up to approx. 10% by weight is only possible with great effort due to the insufficient machinability. The occurrence of long spiral chips in particular causes long machine downtimes, since the chips first have to be removed from the machining area of the machine by hand.

In the case of the alloy according to the invention, on the other hand, the different hard particles serve as chip breakers. The resulting short crumbly chips and/or tangled chips facilitate machinability, which is why the semi-finished products and components from the cast state and the further processed state of the alloy according to the invention have better machinability.

Important exemplary embodiments of the invention are explained using Tables 1 to 12. Cast plates of the copper-nickel-tin alloy according to the invention (example A) and of the reference material R were produced by continuous casting. Furthermore, the continuous casting of pipes with the dimensions (92×72) mm from the exemplary embodiments B and C took place. The chemical composition of the castings can be seen from Table 1.

The exemplary embodiments A to C are characterized by a Ni content of 5.48 to 6.15% by weight, an Sn content of 4.94 to 5.76% by weight, an Fe content of 0.079 to 0. 22% by weight, an Si content of 0.26 to 0.31% by weight, a B content of 0.14 to 0.20% by weight, a P content of 0.048 to 0.072% by weight. -% and characterized by a remainder of copper. The reference material R belongs to the conventional copper-nickel-tin alloys that correspond to the state of the art. It has a Ni content of 5.78% by weight, an Sn content of 5.75% by weight, a P content of about 0.032% by weight and the remainder is copper. <b><u>Table 1:</u></b> Chemical composition of the exemplary embodiments A, B and C and the reference material R (in % by weight) legs Cu no sn feet si B P A rest 6:15 5.76 0.220 0.28 0.14 0.072 B rest 6.06 5.35 0.079 0.26 0.18 0.061 C rest 5.48 4.94 0.200 0.31 0.20 0.048 R rest 5.78 5.75 - - - 0.032

The structure of the continuously cast plates of the reference material R shows gas and shrinkage pores as well as Sn-rich segregations, especially at the grain boundaries.

In contrast to the reference material R, continuous casting has the Embodiments A to C a uniformly solidified, pore-free and segregation-free structure due to the effect of the crystallization nuclei.

The metallic matrix of the as-cast state of the exemplary embodiment A consists of a copper mixed crystal with, based on the overall structure, approx. 10 to 15% by volume of island-shaped embedded first phase components, which can be specified with the empirical formula Cu _h Ni _k Sn _m and a ratio (h+k)/m of the element contents in atomic % from 2 to 6. The compounds CuNi ₁₄ Sn ₂₃ and CuNi ₉ Sn ₂₀ with a ratio (h+k)/m of 3.4 and 4 could be determined. In addition, in the metallic matrix with, based on the overall structure, about 5 to 10% by volume of second phase components are embedded in islands, which can be specified with the molecular formula Cu _p Ni _r Sn _s and a ratio (p + r) / s of Have element contents in atomic percent from 10 to 15. The compounds CuNi ₃ Sn ₈ and CuNi ₄ Sn ₇ with a ratio (p+r)/s of 11.5 and 13.3 were detected. The first and second phase components of the metallic matrix are predominantly crystallized in the area of the crystallization nuclei and encase them.

The analysis of the first-class hard particles in the cast state of exemplary embodiment A gave indications of the compound SiB ₆ as a representative of the Si-containing and B-containing phases, of Ni ₆ Si ₂ B as a representative of the Ni-Si borides, of Ni ₃ B as a Representatives of the Ni borides, FeB as a representative of the Fe borides, Ni ₃ P as a representative of the Ni phosphides, Fe ₂ P as a representative of the Fe phosphides, Ni ₂ Si as a representative of the Ni silicides and Fe -rich particles that are present individually and/or as addition compounds and/or mixed compounds in the structure. In addition, these hard particles are encased by tin and/or the first phase components and/or second phase components of the metallic matrix.

During the casting process of the exemplary embodiments A to C, a substructure formed in the primary casting grains. In the cast structure of exemplary embodiments A to C of the invention, these sub-grains have a grain size of less than 10 μm. As a result of the sub-grain structure and the hard particles precipitated in the structure of the exemplary embodiments A to C of the invention, the hardness HB of the as-cast state of the exemplary embodiments is significantly higher than the hardness of the continuously cast R (Table 2). <b><u>Table 2:</u></b> Hardness HB 2.5/62.5 of the cast condition and the aged condition of the exemplary embodiments A to C and the reference material R legs Continuous casting hardness HB 2.5/62.5 Continuous casting + aging Hardness HB 2.5/62.5 330°C/3h/air 400°C/3h/air 470°C/3h/air A 169 - 173 - B 142 155 158 162 C 156 168 178 180 R 94 - 145 -

Table 2 also shows the hardness values that were determined on the continuously cast alloys A to C and R aged at 330, 400 and 470°C for a period of 3 hours. The increase in hardness from 94 to 145 HB is greatest for the reference material R. This hardening is particularly due to thermally activated segregation of the Sn-rich phase in the structure. The phase components enriched with tin are separated in the structure of the exemplary embodiments A to C in a much finer way in the area of the hard particles. For this reason, the hardness of the alloy A aged at 400° C. increases only slightly from 169 to 173 HB. The hardness HB of the exemplary embodiment C also does not increase so markedly from 156 to 178 as a result of the aging.

A project of the invention is to maintain the good cold workability of conventional copper-nickel-tin alloys despite the

incorporation of hard particles. To check the degree of achievement of this goal, production program 1 was carried out with the continuously cast plates of alloys A and R according to Table 3. This production program consisted of a cycle of cold forming and annealing, with the cold rolling steps being carried out with the maximum possible degree of cold forming.

Due to the high hardness of the cast condition of example A, it was annealed at a temperature of 740° C. for a period of 2 hours and then cooled in water in an accelerated manner. As a result, the properties of the as-cast condition of A and R in terms of strength and hardness were equalized.

The degrees of cold forming ε of 57 and 91% that can be achieved for example A underline that the alloy according to the invention, despite the content of hard particles, can achieve and even surpass the deformation properties of the conventional copper-nickel-tin alloy R.

The temperature sensitivity of the reference material R with regard to the formation of the Sn-rich segregations was also evident during the annealing between the two cold-forming steps (No. 4 in Table 3). For this reason, the annealing temperature of 740°C used for the intermediate anneal of the alloy A cold-rolled plate had to be lowered to 690°C for R. <b><u>Table 3:</u></b> Production program 1 of strips from the continuously cast plates of example A and the reference material R No. manufacturing steps 1 Continuous cast plates of alloys A and R 2 Annealing of the cast plate of the Leg. A: 740°C/2h+water quench 3 cold rolling legs A: from 11 to 4.70 mm (ε= 57%, ϕ= 0.8) legs R: from 24.5 to 12.1 mm (ε= 50%, ϕ= 0.7) 4 glow legs A: 740°C/2h + water quench legs R: 690°C/2h + water quench 5 cold rolling legs A: from 4.70 to 0.4 mm (ε= 91%, ϕ= 2.4) legs R: from 12.1 to 2.33 mm (ε= 81%, ϕ= 1.6) 6 Aging: 300°C/4h, 400°C/3h, 450°C/3h + air cooling

After the execution of production program 1, the characteristic values of the strips of materials A and R were determined after the last cold rolling and after aging, which are listed in Table 4.

It is clear that the strength and hardness of the cold-rolled strips and those stripped at 300°C from example A are higher than the respective properties of the strips of reference material R.

Favored by the high content of hard particles, a recrystallization of the structure of alloy A takes place from a temperature of approx. 400°C. This recrystallization leads to a drop in strength and hardness, so that the effect of precipitation hardening and spinodal segregation cannot come into play. Since no recrystallization of the structure can be observed with the reference material R up to 450°C, the values for R _m , R _p0.2 and for the hardness, especially after aging at 400° C., are higher at R than in exemplary embodiment A.

After aging at 450° C., the structure of the further processed exemplary embodiment A contains the hard particles of the second class (in 3 labeled 3).

Furthermore, further phases have separated out in the structure of the further processed Alloy A. These include the in 3 continuous precipitates of the (Cu,Ni)-Sn system denoted by 4 and the hard particles of the third class.

= noch nicht vollständig rekristallisiert Leg. Auslagerung [°C/ h] Korngröße [µm] Elektrische Leitfähigkeit [%IACS] R_m [MPa] R_p0,2 [MPa] A [%] E [GPa] Härte HV1 A - - 10,6 974 917 3,2 119 311 300°C/4h - 15,6 969 927 5,3 125 320 400°C/3h

< 2 23,2 704 679 17,5 126 237 450°C/3h < 1 24,5 591 575 22,3 126 193 R - - 10,7 838 787 7,2 120 267 300°C/4h - 13,8 910 874 9,2 118 297 400°C/3h - 22,0 793 735 13,6 108 264 450°C/3h - 23,2 610 508 23,0 124 195 The size of the third-class hard particles of less than 3 μm is characteristic of the further processed alloy according to the invention. For the further processed exemplary embodiment A of the invention, after aging at 450° C., it is even less than 1 μm (in 4 marked 5). <b><u>Table 4:</u></b> Grain size, electrical conductivity and mechanical characteristics of the cold-rolled and aged strips of alloys A and R after running through production program 1 (Table 3)

= not yet fully recrystallized legs Aging [°C/h] grain size [µm] Electrical conductivity [%IACS] R _m [MPa] R _p0.2 [MPa] A [%] E [GPa] Hardness HV1 A - - 10.6 974 917 3.2 119 311 300°C/4h - 15.6 969 927 5.3 125 320 400°C/3h

< 2 23.2 704 679 17.5 126 237 450°C/3h < 1 24.5 591 575 22.3 126 193 R - - 10.7 838 787 7.2 120 267 300°C/4h - 13.8 910 874 9.2 118 297 400°C/3h - 22.0 793 735 13.6 108 264 450°C/3h - 23.2 610 508 23.0 124 195

In order to reduce the influence of cold workability and the recrystallization temperature on the properties of the individual alloys, another production program was carried out. This production program 2 pursued the goal of processing the continuously cast plates of materials A and R into strips by means of cold forming and annealing, with identical parameters for the degrees of cold forming and the annealing temperatures being used in each case (Table 5). Due to the high hardness of the cast condition of example A, this was again annealed before the first cold-rolling step at a temperature of 740° C. for a period of 2 hours and then cooled in water in an accelerated manner. As with production program 1, this resulted in the properties of the cast condition of A and R being matched with regard to strength and hardness. <b><u>Table 5:</u></b> Production program 2 of strips from the continuously cast plates of example A and the reference material R No. manufacturing steps 1 Continuous cast plates of alloys A and R 2 Annealing of the cast plate of the Leg. A: 740°C/2h+water quench 3 Cold rolling: from 9 to 6 mm (ε= 33%, ϕ= 0.4) 4 Annealing: 690°C/2h + water quenching 5 Cold rolling: from 6 to 3.5 mm (ε= 42%, ϕ= 0.5) 6 Annealing: 690°C/1h + water quenching 7 Cold rolling: from 3.5 to 3.0 mm (ε= 14%, ϕ= 0.15) 8th Aging: 400°C/3h, 450°C/3h, 500°C/3h + air cooling

After the last cold-rolling step to the final thickness of 3.0 mm, the alloy A strip exhibited the highest strength and hardness values (Table 6).

Due to the three-hour aging at 400°C, the increase in strength R _m (from 498 to 717 MPa) and R _p0.2 (from 439 to 649 MPa) as well as the hardness HB (from 166 to 230 MPa) falls due to the spinodal segregation of the structure ) is most evident in alloy R. However, the structure of the aged states of alloy R is very uneven with a grain size between 5 and 30 µm. In addition, the structure of the aged states of the reference material R is characterized by discontinuous precipitations of the system (Cu, Ni)-Sn (in 1 and 2 denoted by 1). The structure of the further processed state of the reference material R also contains Ni phosphides (in 1 and 2 labeled 2).

In contrast, the structure of the aged strips of embodiment A of the invention is very uniform with a grain size of 2 to 8 μm. In addition, the structure of Example A lacks the discontinuous precipitates even after aging at 450°C for 3 hours followed by air cooling. In contrast, second-class hard particles can be detected in the structure. These phases are in 5 and 6 marked with 3.

Furthermore, further phases have separated out in the structure of the further processed Alloy A. These include the in figure 5 continuous precipitates of the system (Cu,Ni)-Sn denoted by 4 and the hard particles of the third class. For the further processed exemplary embodiment A of the invention, the size of the third-class hard particles after aging at 450° C. is even less than 1 μm (in 6 marked 5).

= ungleichmäßig Leg. Auslagerung [°C/ h] Korngröße [µm] Elektrische Leitfähigkeit [%IACS] R_m [MPa] R_p0,2 [MPa] A [%] E [GPa] Härte HBW 1/30 A - - 11,6 556 498 25,1 113 188 400°C/3h 2-8 15,1 690 618 21,4 132 222 450°C/3h 2-8 16,8 666 534 22,1 126 211 500°C/3h 2-8 16,7 614 444 24,4 124 190 R - - 11,2 498 439 27,9 104 166 400°C/3h

5-30 15,2 717 649 17,8 132 230 450°C/3h

5-30 17,0 705 591 20,6 121 219 500°C/3h

5-20 18,6 628 420 24,6 118 190 After aging at 400°C/3h/air, the strengths R _m and R _p0.2 of the strips of alloy A assume the values of 690 and 618 MPa as a result of the spinodal segregation of the structure. R _m and R _p0.2 are therefore lower than the characteristic values of the correspondingly aged state of alloy R. This is due to the fact that in exemplary embodiment A the Ni content, which is bound in the hard particles, is responsible for the strength-increasing spinodal segregation of the structure is missing. If the strength level of R is required, it is possible to add a higher proportion of the alloying element nickel to the alloy according to the invention. <b><u>Table 6:</u></b> Grain size, electrical conductivity and mechanical characteristics of the cold-rolled and aged strips of alloys A and R after running through production program 2 (Table 5)

= uneven legs Aging [°C/h] grain size [µm] Electrical conductivity [%IACS] R _m [MPa] R _p0.2 [MPa] A [%] E [GPa] Hardness HBW 1/30 A - - 11.6 556 498 25.1 113 188 400°C/3h 2-8 15.1 690 618 21:4 132 222 450°C/3h 2-8 16.8 666 534 22:1 126 211 500°C/3h 2-8 16.7 614 444 24.4 124 190 R - - 11.2 498 439 27.9 104 166 400°C/3h

5-30 15.2 717 649 17.8 132 230 450°C/3h

5-30 17.0 705 591 20.6 121 219 500°C/3h

5-20 18.6 628 420 24.6 118 190

The next step included testing the hot formability of the continuously cast alloys A and R. For this purpose, the cast plates were hot rolled at a temperature of 720°C (Table 7). The parameters of production program 2 were adopted for the further process steps of cold forming and intermediate annealing. <b><u>Table 7:</u></b> Production program 3 of strips from the continuously cast plates of example A and the reference material R No. manufacturing steps 1 Continuous cast plates of alloys A and R 2 legs A, R: Hot rolling at 720°C + water quenching 3 cold rolling alloy A: from 9 to 6 mm (ε= 33%, ϕ= 0.4) 4 glow leg A: 690°C/2h + water quench 5 cold rolling alloy A: from 6 to 3.5 mm (ε= 42%, ϕ= 0.5) 6 glow leg A: 690°C/1h + water quench 7 cold rolling alloy A: from 3.5 to 3.0 mm (ε= 14%, ϕ= 0.15) 8th Outsourcing Leg. A: 400°C/3h, 450°C/3h + air cooling

During the hot rolling of the cast plates of the reference alloy R, deep hot cracks formed after just a few passes, which led to the plates failing due to fracture.

In contrast, the cast plates of embodiment A of the invention could be hot-rolled without damage and, after several cold-rolling and annealing processes, could be manufactured to the final thickness of 3.0 mm. The properties of the outsourced strips (Tab. 8) largely correspond to those of the strips that were produced without hot forming using production program 2 (Tab. 6).

The microstructure of the strips from embodiment A of the alloy according to the invention, which were produced with and without a hot forming step, is also comparable. That's how it goes 7 and 8 the uniform structure of the strips from example A emerges, which were produced with a hot forming stage and a final aging at 400° C./3 h/air cooling. In 7 and 8 the hard particles of the second class, labeled 3, can again be seen.

Continue going out 7 the continuous precipitations of the system (Cu,Ni)-Sn, denoted by 4, as well as the hard particles of the third class. In the structure of the further processed variant of example A, the hard particles of the third class even assume a size of less than 1 µm (with 5 in 8 designated).

The analysis of the hard particles of the second and third class in this further processed state of example A gave indications of the compound SiB ₆ as a representative of the Si-containing and B-containing phases, of Ni ₆ Si ₂ B as a representative of the Ni-Si borides Ni ₃ B as a representative of Ni borides, on FeB as a representative of Fe borides, on Ni ₃ P as a representative of Ni phosphides, on Fe ₂ P as a representative of Fe phosphides, on Ni ₂ Si as a representative of Ni Silicides and Fe-rich particles that are present individually and as addition compounds and/or mixed compounds in the structure. In addition, these hard particles are coated with precipitates of the (Cu, Ni)-Sn system. <b><u>Table 8:</u></b> Grain size, electrical conductivity and mechanical characteristics of the cold-rolled and aged strips of example A after running through production program 3 (table 7) legs Aging [°C/h] grain size [µm] Electrical conductivity [%IACS] Rm [MPa] R _p0.2 [MPa] A [%] E [GPa] Hardness HBW 1/30 A - - 11.8 554 501 23.8 110 185 400°C/3h 3-10 15.3 679 610 21.0 127 217 450°C/3h 3-10 16.8 658 535 20.8 126 205

The subsequent experimental stage included testing the hot forming behavior of embodiment A of the invention at the higher hot rolling temperature of 780°C. The aim was also to reduce the number of cold rolling/annealing cycles in production program 3. This measure made it possible to investigate the cold formability of the hot-rolled strip condition of alloy A. The individual process steps of production program 4 are shown in Table 9. <b><u>Table 9:</u></b> Production program 4 of strips from the continuously cast plates of example A No. manufacturing steps 1 Alloy A continuous cast slabs 2 legs A: Hot rolling at 780°C + water quenching 3 cold rolling alloy A: from 9 to 1.4 mm (ε= 84%, ϕ= 1.9) 4 Glow Leq. A: 690°C/1h + water quench 5 cold rolling Leq. A: from 1.4 to 1.2 mm (ε= 14%, ϕ= 0.15) 6 Aging: 350°C/3h, 400°C/3h, 450°C/3h, 500°C/3h + air cooling

Even at the higher hot rolling temperature, Alloy A continuously cast slabs showed excellent hot formability. The hot-rolled plates could also be cold-rolled without any problems with an extremely high cold-forming degree ε of 84%. In order to make the outsourcing result comparable with the result of the previous production program 3 can, the last cold rolling step took place after recrystallization annealing at 690°C with the same degree of cold deformation ε of 14%.

After the strips have been aged in the temperature range of 350 to 500°C, the grain size of the very uniform structure is 5 to 10 µm (Table 10). In particular at the aging temperature of 400° C., the spinodal segregation of the structure of the alloy according to the invention leads to a pronounced increase in strength and hardness. The tensile strength R _m increases from 557 MPa in the cold-rolled condition to 692 MPa in the aged condition. The HB hardness also increases from 177 to 210. <b><u>Table 10:</u></b> Grain size, electrical conductivity and mechanical characteristics of the cold-rolled and aged strips of Alloy A after running through production program 4 (Table 9) legs Aging [°C/h] grain size [µm] Electrical conductivity [%IACS] R _m [MPa] R _p0.2 [MPa] A [%] E [GPa] Hardness HBW 1/30 A - - 11.6 557 520 22.2 - 177 350°C/3h 5-10 13.8 674 607 23.3 143 204 400°C/3h 5-10 15.2 692 614 20.1 150 210 450°C/3h 5-10 17.2 659 519 21:9 128 193 500°C/3h 5-10 15.8 598 437 25.0 128 170

In plant, device, engine and mechanical engineering, components with larger dimensions are required for numerous applications. For example, this is often the case in the field of plain bearings. The production of the corresponding components requires a starting material of correspondingly large formats. Due to the limited manufacturability of castings of any size, there is therefore a need to adjust the required material properties as far as possible by means of small degrees of cold forming.

Table 11 shows the ones used in production program 5 Process steps listed. The production was carried out with a cycle of cold forming and annealing. Again, only the Alloy A cast slabs were annealed at 740°C prior to the first cold rolling.

The first cold rolling of the Alloy R cast slab and the annealed Alloy A cast slab was carried out with a deformation ε of 16%. After annealing at 690° C., cold rolling took place with ε of 12%. Finally, the strips were aged at temperatures of 350, 400 and 450°C. <b><u>Table 11:</u></b> Production program 5 of strips from the continuously cast slabs of the exemplary embodiment A and the reference material R No. manufacturing steps 1 Continuous cast plates of alloys A and R 2 Annealing of the cast plates of the Leg. A: 740°C/2h+water quench 3 cold rolling alloy A, R: from 9 to 7.6 mm (ε= 16%, ϕ= 0.17) 4 glow leg A, R: 690°C/2h + water quench 5 cold rolling alloy A, R: from 7.6 to 6.7 mm (ε= 12%, ϕ= 0.126) 6 Aging: 350°C/3h, 400°C/3h, 450°C/3h + air cooling

The low cold forming of the first cold rolling step of ε= 16% was not enough to eliminate the dendritic and coarse-grained structure of the reference material R together with the subsequent annealing at 690°C. In addition, this thermomechanical treatment increased the coverage of the grain boundaries of alloy R with Sn-rich segregations.

During the second cold rolling step, cracks developed along the dendritic structure and along the grain boundaries of R, which are covered with Sn-rich segregations. These cracks run from the surface deep into the interior of the strip.

The crack-free and uniform structure of the belts of example A is characterized by the arrangement of the hard particles of the second and third class. As with the previous production programs, the hard particles of the third class also have a size of less than 1 μm after this production program 5.

The resulting strip properties after the last cold rolling and after aging are shown in Table 12. Due to the high density of cracks, it was not possible to take damage-free tensile specimens from the ribbons of material R. Thus, only the metallographic examination and the hardness measurement could be carried out on these strips.

Exemplary embodiment A has a high degree of aging ability, which is expressed by interaction of the mechanisms of precipitation hardening and spinodal segregation of the structure. For example, the characteristic values R _m and R _{p0.2 increase} from 518 to 633 and from 451 to 575 MPa as a result of aging at 400°C. <b><u>Table 12:</u></b> Grain size, electrical conductivity and mechanical characteristics of the cold-rolled and aged strips of alloys A and R after running through production program 5 (table 11) ■ = dendritic, with Sn-rich segregations legs outsourcing grain size [µm] Electrical conductivity [%IACS] Rm [MPa] R _p0.2 [MPa] A [%] E [GPa] Hardness HBW 1/30 A - - 11.6 518 451 21:8 124 188 350°C/3h 20 13.4 610 530 23.5 130 212 400°C/3h 20-25 14.4 633 575 19.1 116 217 450°C/3h 20-25 16.0 630 496 17.4 108 204 R - ■ - Not possible due to crack formation! 175 350°C/3h ■ - 242 400°C/3h ■ - 229 450°C/3h ■ - 217

As a result, it can be stated that the degree of precipitation hardening and the degree of spinodal segregation of the structure of the invention can be adapted to the required material properties by varying the chemical composition, the degree of deformation for the cold forming(s) and by varying the aging conditions. In this way, it is possible to tailor the strength, hardness, ductility and electrical conductivity of the alloy according to the invention to the intended area of use.

Reference List

1

Discontinuous precipitations of the (Cu,Ni)-Sn system

2

Ni phosphides

3

Second class hard particles

4

Continuous precipitations of the system (Cu, Ni)-Sn as well as third-class hard particles

5

Third class hard particles

Claims (18)

1. High-strength copper/nickel/tin alloy in the as-cast state having excellent castability, hot-formability and cold-formability, high resistance against abrasive wear, adhesive wear and fretting wear and improved corrosion resistance and stress relaxation resistance, comprising (in % by weight):

   from 2.0 to 10.0% Ni,

   from 2.0 to 10.0% Sn,

   from 0.01 to 1.5% Si,

   from 0.01 to 1.0% Fe,

   from 0.002 to 0.45% B,

   from 0.001 to 0.15% P,

   optionally also up to a maximum of 2.0% Co,

   optionally also up to a maximum of 2.0% Zn,

   optionally also up to a maximum of 0.25% Pb,

   the balance being copper and inevitable impurities,

   characterised in that

   - the ratio Si/B of the element contents in % by weight of the elements silicon and boron is a minimum of 0.4 and a maximum of 8;

   - in that after the casting the following structural components are present in the alloy:

   a) an Si-containing and P-containing metal matrix having, relative to the overall structure,

   a1) up to 30% by volume of first phase components which with the molecular formula Cu_hNi_kSn_m have a ratio (h+k)/m of the element contents in atom. % of from 2 to 6,

   a2) up to 20% by volume of second phase components which with the molecular formula Cu_pNi_rSn_s have a ratio (p+r)/s of the element contents in atom. % of from 10 to 15, and

   a3) with a balance of mixed copper crystal;

   b) phases which are contained in the structure relative to the overall structure,

   b1) at from 0.01 to 10% by volume as Si-containing and B-containing phases which are formed as silicon borides and as boron silicates and/or as boron phosphorus silicates,

   b2) at from 1 to 15% by volume as Ni/Si borides with the molecular formula Ni_xSi₂B with x = 4 to 6,

   b3) at from 1 to 15% by volume as Ni borides,

   b4) at from 0.1 to 5% by volume as Fe borides,

   b5) at from 1 to 5% by volume as Ni phosphides,

   b6) at from 0.1 to 5% by volume as Fe phosphides,

   b7) at from 1 to 5% by volume as Ni silicides,

   b8) at from 0.1 to 5% by volume as Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or mixed compounds and which are surrounded by tin and/or the first phase components and/or

   the second phase components;

   - in that during casting the Si-containing and B-containing phases which are in the form of silicon borides, the Ni-Si borides, the Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and the Fe silicides and/or Fe-rich particles which are present individually and/or as addition compounds and/or mixed compounds constitute germs for a uniform crystallisation during the solidification/cooling of the melt so that the first phase components and/or the second phase components are uniformly distributed in an island-like and/or network-like manner in the structure;

   - in that the Si-containing and B-containing phases which are in the form of boron silicates and/or boron phosphorus silicates together with the phosphorus silicates act as a wear-prevention and corrosion-prevention coating on the semifinished products and components of the alloy.
2. High-strength copper/nickel/tin alloy after annealing or after hot-forming and/or cold-forming in addition to annealing with excellent castability, hot-formability and cold-formability, high resistance against abrasive wear, adhesive wear and fretting wear and improved corrosion resistance and stress relaxation resistance, comprising (in % by weight):

   from 2.0 to 10.0% Ni,

   from 2.0 to 10.0% Sn,

   from 0.01 to 1.5% Si,

   from 0.01 to 1.0% Fe,

   from 0.002 to 0.45% B,

   from 0.001 to 0.15% P,

   optionally also up to a maximum of 2.0% Co,

   optionally also up to a maximum of 2.0% Zn,

   optionally also up to a maximum of 0.25% Pb,

   the balance being copper and inevitable impurities,

   characterised in that

   - the ratio Si/B of the element contents in % by weight of the elements silicon and boron is a minimum of 0.4 and a maximum of 8;

   - in that, after the further processing of the alloy by means of at least one annealing operation or by means of at least one hot-forming and/or cold-forming in addition to at least one annealing operation, the following structural components are present:

   A) a metal matrix having, relative to the overall structure,

   A1) up to 15% by volume of first phase components which with the molecular formula Cu_hNi_kSn_m have a ratio (h+k)/m of the element contents in atom. % of from 2 to 6,

   A2) up to 10% by volume of second phase components which with the molecular formula Cu_pNi_rSn_s have a ratio (p+r)/s of the element contents in atom. % of from 10 to 15, and

   A3) a balance of mixed copper crystal;

   B) phases which, relative to the overall structure, are contained in the structure

   B1) at from 2 to 35% by volume as Si-containing and B-containing phases which are in the form of silicon borides and boron silicates and/or boron phosphorus silicates, Ni-Si borides with the molecular formula Ni_xSi₂B with x = 4 to 6,

   Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles, which are present individually and/or as addition compounds and/or

   mixed compounds and are coated with precipitations of the system (Cu, Ni)-Sn,

   B2) at up to 80% by volume as continuous precipitations of the system (Cu, Ni)-Sn,

   B3) are contained in the structure at from 2 to 35% by volume as Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles which are present individually and/or as addition compounds and/or mixed compounds, are coated with precipitations of the system (Cu,

   Ni)-Sn and have a size of less than 3 µm;

   - in that the Si-containing and B-containing phases which are in the form of silicon borides, the Ni-Si borides, Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and the Fe silicides and/or Fe-rich particles which are present individually and/or as addition compounds and/or mixed compounds constitute germs for a static and dynamic recrystallisation of the structure during the further processing of the alloy, whereby a uniform and fine-grained structure is produced;

   - in that the Si-containing and B-containing phases, which are in the form of boron silicates and/or boron phosphorus silicates together with the phosphorus silicates act as a wear-prevention and corrosion-prevention coating on the semifinished products and components of the alloy.
3. Copper/nickel/tin alloy according to claim 1 or 2, characterised in that the elements nickel and tin are each contained at from 3.0 to 9.0%.
4. Copper/nickel/tin alloy according to any one of claims 1 to 3, characterised in that the element silicon is contained at from 0.05 to 0.9%.
5. Copper/nickel/tin alloy according to any one of claims 1 to 4, characterised in that the element iron is contained at from 0.02 to 0.6%.
6. Copper/nickel/tin alloy according to any one of claims 1 to 5, characterised in that the element boron is contained at from 0.01 to 0.4%.
7. Copper/nickel/tin alloy according to any one of claims 1 to 6, characterised in that the element phosphorus is contained at from 0.01 to 0.15%.
8. Copper/nickel/tin alloy according to any one of claims 1 to 7, characterised in that the alloy is free from lead with the exception of any inevitable impurities.
9. Method for producing end products or components with a form close to the end product from a copper/nickel/tin alloy according to any one of claims 1 to 8 using the sand casting method, shell mould casting method, precision casting method, full mould casting method, die-casting method or lost-foam method.
10. Method for producing strips, metal sheets, plates, bolts, round wires, profiled wires, round rods, profiled rods, hollow rods, pipes and profiles from a copper/nickel/tin alloy according to any one of claims 1 to 8, using the chilled mould casting method or the continuous or semi-continuous strand casting method.
11. Method according to claim 10, characterised in that the further processing of the casting state involves carrying out at least one hot-forming operation in the temperature range from 600 to 880°C.
12. Method according to any one of claims 9 to 11, characterised in that at least one annealing processing operation is carried out in the temperature range from 170 to 880°C with a duration of from 10 minutes to 6 hours.
13. Method according to any one of claims 10 to 12, characterised in that the further processing of the as-cast state or the hot-formed state or the annealed as-cast state or the annealed hot-formed state involves carrying out at least one cold-forming operation.
14. Method according to claim 13, characterised in that at least one annealing processing operation is carried out in the temperature range from 170 to 880°C with a duration of from 10 minutes to 6 hours.
15. Method according to either claim 13 or claim 14, characterised in that a stress-relieving annealing operation or precipitation annealing operation is carried out in the temperature range from 170 to 550°C with a duration of from 0.5 to 8 hours.
16. Use of the copper/nickel/zinc alloy according to any one of claims 1 to 8 for gibs or sliding rails, for friction rings or friction plates, for sliding elements or guiding elements in internal combustion engines, valves, turbochargers, gear mechanisms, exhaust gas post-treatment systems, lever systems, brake systems or joint systems, hydraulic units or in machines or installations of general mechanical engineering.
17. Use of the copper/nickel/tin alloy according to any one of claims 1 to 8 for structural elements, conduction elements, guide elements or connection elements in electronics or electrical engineering.
18. Use of the copper/nickel/tin alloy according to any one of claims 1 to 8 for metal objects in the breeding of organisms living in seawater, for percussion instruments, for propellors, wings, ship propellors or hubs for ship building, for housings of water pumps, oil pumps or fuel pumps, for guide wheels, impellers or paddle wheels for pumps or water turbines, for gears, worm gears, helical gears, pressure nuts or spindle nuts and for pipes, seals or connection bolts in the maritime or chemical industry.

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