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

Abstract

A copper-nickel-tin alloy with excellent castability, hot and cold workability, high resistance to abrasive wear, adhesive wear and fretting wear and improved resistance to corrosion and stress relaxation stability, consisting of (in weight %): 2.0-10.0% Ni, 2.0-10.0% Sn, 0.01-1.5% Si, 0.01-1.0% Fe, 0.002-0.45% B, 0.001-0.15% P, selectively up to a maximum of 2.0% Co, optionally also up to a maximum 2.0% Zn, selectively up to a maximum of 0.25% Pb, the residue being copper and unavoidable impurities. The ratio Si/B of the element contents in wt. % of the elements silicon and boron is a minimum 0.4 and a maximum 8 such that the copper-nickel-tin alloy has Si-containing and B-containing phases, phases of the systems Ni—Si—B, Ni—B, Fe—B, Ni—P, Fe—P, Ni—Si, and other Fe-containing phases, which improve the processing and use properties of the alloy.

Claims

1. A copper-nickel-tin alloy 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, the balance being copper and unavoidable impurities, wherein 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; the following microstructure constituents are present in the alloy after casting: a) a Si-containing and P-containing metallic base composition having, based on the overall microstructure, a1) up to 30% by volume of first phase constituents that can be reported by the empirical formula Cu.sub.hNi.sub.kSn.sub.m and have an (h+k)/m ratio of the element contents in atom % of 2 to 6, a2) up to 20% by volume of second phase constituents that can be reported by the empirical formula Cu.sub.pNi.sub.rSn.sub.s and have a (p+r)/s ratio of the element contents in atom % of 10 to 15 and a3) a balance of a solid copper solution; b) phases which, based on the overall microstructure, are present b1) at 0.01% to 10% by volume as Si-containing and B-containing phases, b2) at 1% to 15% by volume as Ni—Si borides having the empirical formula Ni.sub.xSi.sub.2B with x=4 to 6, b3) at 1% to 15% by volume as Ni borides, b4) at 0.1% to 5% by volume as Fe borides, b5) at 1% to 5% by volume as Ni phosphides, b6) at 0.1% to 5% by volume as Fe phosphides, b7) at 1% to 5% by volume as Ni silicides, and b8) at 0.1% to 5% by volume as Fe silicides and/or Fe-rich particles in the microstructure, which are present individually and/or as addition compounds and/or mixed compounds and are ensheathed by tin and/or the first phase constituents and/or the second phase constituents; in the course of casting the Si-containing and B-containing phases 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 that are present individually and/or as addition compounds and/or mixed compounds constitute seeds for uniform crystallization during the solidification/cooling of the melt, such that the first phase constituents and/or the second phase constituents are distributed uniformly in the microstructure in the form of islands and/or in the form of a mesh; the Si-containing and B-containing phases that are in the form of boron silicates and/or boron phosphorus silicates, together with phosphorus silicates, assume the role of a wear-protecting and corrosion-protecting coating on semifinished materials and components of the alloy.

2. The copper-nickel-tin alloy as claimed in claim 1, wherein the elements nickel and tin are each present at 3.0% to 9.0%.

3. The copper-nickel-tin alloy as claimed in claim 1, wherein the element silicon is present at 0.05% to 0.9%.

4. The copper-nickel-tin alloy as claimed in claim 1, wherein the element iron is present at 0.02% to 0.6%.

5. The copper-nickel-tin alloy as claimed in claim 1, wherein the element boron is present at 0.01% to 0.4%.

6. The copper-nickel-tin alloy as claimed in claim 1, wherein the element phosphorus is present at 0.01% to 0.15%.

7. The copper-nickel-tin alloy as claimed in claim 1, wherein the alloy is free of lead apart from any unavoidable impurities.

8. A copper-nickel-tin alloy 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, the balance being copper and unavoidable impurities, wherein 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; after a further processing of the alloy by at least one annealing operation or by at least one hot forming operation and/or cold forming operation, as well as at least one annealing operation, the following microstructure constituents are present: A) a metallic base composition having, based on the overall microstructure, A1) up to 15% by volume of first phase constituents that can be reported by the empirical formula Cu.sub.hNi.sub.kSn.sub.m, and have an (h+k)/m ratio of the element contents in atom % of 2 to 6, A2) up to 10% by volume of second phase constituents that can be reported by the empirical formula Cu.sub.pNi.sub.rSn.sub.s and have a (p+r)/s ratio of the element contents in atom % of 10 to 15 and A3) a balance of a solid copper solution; B) phases which, based on the overall microstructure, are present B1) at 2% to 35% by volume as Si-containing and B-containing phases, Ni—Si borides having the empirical formula Ni.sub.xSi.sub.2B with x=4 to 6, Ni borides, Fe borides, Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles in the microstructure, which are present individually and/or as addition compounds and/or mixed compounds and are ensheathed by precipitates of the (Cu, Ni)—Sn system, B2) at up to 80% by volume as continuous precipitates of the (Cu, Ni)—Sn system in the microstructure, B3) at 2% to 35% by volume as Ni phosphides, Fe phosphides, Ni silicides and as Fe silicides and/or Fe-rich particles in the microstructure that are present individually and/or as addition compounds and/or mixed compounds, are ensheathed by precipitates of the (Cu, Ni)—Sn system and have a size of less than 3 μm; the Si-containing and B-containing phases that 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 that are present individually and/or as addition compounds and/or mixed compounds constitute seeds for static and dynamic recrystallization of the microstructure during the further processing of the alloy, which enables the establishment of a uniform and fine-grain microstructure; the Si-containing and B-containing phases that are in the form of boron silicates and/or boron phosphorus silicates, together with phosphorus silicates, assume the role of a wear-protecting and corrosion-protecting coating on semifinished materials and components of the alloy.

9. The copper-nickel-tin alloy as claimed in claim 8, wherein the elements nickel and tin are each present at 3.0% to 9.0%.

10. The copper-nickel-tin alloy as claimed in claim 8, wherein the element silicon is present at 0.05% to 0.9%.

11. The copper-nickel-tin alloy as claimed in claim 8, wherein the element iron is present at 0.02% to 0.6%.

12. The copper-nickel-tin alloy as claimed in claim 8, wherein the element boron is present at 0.01% to 0.4%.

13. The copper-nickel-tin alloy as claimed in claim 8, wherein the element phosphorus is present at 0.01% to 0.15%.

14. The copper-nickel-tin alloy as claimed in claim wherein the alloy is free of lead apart from any unavoidable impurities.

Description

(1) Examples of the invention are explained in more detail below that include references to the drawings, in which:

(2) FIG. 1 and FIG. 2 show discontinuous precipitates of the (Cu, Ni)-Sn system and Ni phosphides in the microstructure of an age-hardened reference material R.

(3) FIG. 3 shows hard particles of the second class and continuos precipitates of the (Cu, Ni)-Sn system in the microstructure of working example A.

(4) FIG. 4 shows hard particles of the third class in the microstructure of working example A.

(5) FIG. 5 shows hard particles of the second class and continuous precipitates of the (Cu, Ni)-Sn system in the microstructure of working example A.

(6) FIG. 6 shows hard particles of the second class and hard particles of the third class in the microstructure of working example A.

(7) FIG. 7 shows hard particles of the second class and continuous precipitates of the (Cu, Ni)-Sn system in the microstructure of working example A.

(8) FIG. 8 shows hard particles of the second class and hard particles of the third class in the microstructure of a further-processed variant of working example A.

(9) Important working examples of the invention are illustrated by Tables 1 to 12. Cast plates of the copper-nickel-tin alloy of the invention (working example A) and of the reference material R were produced by strand casting. Furthermore, pipes of dimensions (92×72) mm were strand-cast from working examples B and C. The chemical composition of the casts is apparent from Table 1.

(10) The working examples A-C are characterized by a Ni content of 5.48% to 6.15% by weight, a Sn content of 4.94% to 5.76% by weight, a Fe content of 0.079% to 0.22% by weight, a 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 by a balance of copper. The reference material R is one of the conventional copper-nickel-tin alloys which correspond to the prior art. It has a Ni content of 5.78% by weight, a Sn content of 5.75% by weight, a P content of about 0.032% by weight, and a balance of copper.

(11) TABLE-US-00001 TABLE 1 Chemical composition of working examples A, B and C and of the reference material R (in % by weight) Alloy Cu Ni Sn Fe Si B P A Balance 6.15 5.76 0.220 0.28 0.14 0.072 B Balance 6.06 5.35 0.079 0.26 0.18 0.061 C Balance 5.48 4.94 0.200 0.31 0.20 0.048 R Balance 5.78 5.75 — — — 0.032

(12) The microstructure of the strand-cast plates of the reference material R has gas pores, shrinkage pores, and Sn-rich segregations particularly at the grain boundaries.

(13) By contrast with the reference material R, the strand casting of the working examples A to C, due to the effect of the crystallization seeds, has a uniformly solidified, pore-free and segregation-free microstructure.

(14) The metallic base material of the cast state of the working example A consists of a solid copper solution with, based on the overall microstructure, about 10% to 15% by volume of intercalated first phase constituents in the form of islands, which can be reported by the empirical formula Cu.sub.hNi.sub.kSn.sub.m and have a ratio (h+k)/m of the element contents in an atomic % of 2 to 6. It was possible to detect the compounds CuNi.sub.14Sn.sub.23 and CuNi.sub.9Sn.sub.20 with a ratio (h+k)/m of 3.4 and 4. Also, second phase constituents that can be reported by the empirical formula Cu.sub.pNi.sub.rSn.sub.s, and have a ratio (p+r)/s of the element contents in atom % of 10 to 15, are intercalated in the form of islands in the metallic base material at about 5% to 10% by volume based on the overall microstructure. The compounds CuNi.sub.3Sn.sub.8 and CuNi.sub.4Sn.sub.7 were detected with a ratio (p+r)/s of 11.5 and 13.3. The first and second phase constituents of the metallic base material are predominantly crystallized in the region of the crystallization seeds and ensheath them.

(15) The analysis of the hard particles of the first class in the cast state of the working example A revealed the compound SiB.sub.6 as a representative of the Si-containing and B-containing phases, Ni.sub.6Si.sub.2B as a representative of the Ni—Si borides, Ni.sub.3B as a representative of the Ni borides, FeB as a representative of the Fe borides, Ni.sub.3P as a representative of the Ni phosphides, Fe.sub.2P as a representative of the Fe phosphides, Ni.sub.2Si as a representative of the Ni silicides, and Fe-rich particles, which are present in the microstructure individually and/or as addition compounds and/or mixed compounds. In addition, these hard particles are ensheathed by tin and/or the first phase constituents and/or second phase constituents of the metallic base material.

(16) During the process of casting the working examples A to C, a substructure formed in the primary cast grains. These subgrains in the cast microstructure of the working examples A to C of the invention have a grain size of less than 10 μm. As a result of the subgrain structure and the hard particles precipitated in the microstructure of the working examples A to C of the invention, the hardness HB of the cast state of the working examples is well above the hardness of the strand casting of the reference material R (Table 2).

(17) TABLE-US-00002 TABLE 2 Hardness HB 2.5/62.5 of the cast state and of the state of working examples A to C and of reference material R that have been age-hardened Strand casting Strand casting + age hardening Hardness HB Hardness HB 2.5/62.5 Alloy 2.5/62.5 330° C./3 h/air 400° C./3 h/air 470° C./3 h/air A 169 — 173 — B 142 155 158 162 C 156 168 178 180 R 94 — 145 —

(18) Likewise shown in Table 2 are the hardness values that have been ascertained on the strand casting of alloys A to C and R that has been age-hardened at 330, 400 and 470° C. for a duration of 3 hours. The rise in hardness from 94 to 145 HB is at its greatest for the reference material R. This hardening is particularly attributable to the thermally activated formation of the segregation of the Sn-rich phase in the microstructure. The tin-enriched phase constituents precipitate out in much finer form in the region of the hard particles in the microstructure of the working examples A to C. For this reason, the hardness of the state of the alloy A after age hardening at 400° C. rose only slightly from 169 to 173 HB. The rise in the hardness HB of the working example C from 156 to 178 as a result of the age hardening is also not as marked.

(19) One intention of the invention is that of maintaining the good cold formability of the conventional copper-nickel-tin alloys in spite of the introduction of hard particles. To verify the degree to which this aim is achieved, the manufacturing program 1 with the strand-cast plates of the alloys A and R according to Table 3 was conducted. This manufacturing program consisted of one cycle of cold forming and annealing operations, wherein the cold rolling steps were each carried out with the maximum possible degree of cold forming.

(20) Due to the high hardness of the cast state of working example A, it was calcined at the temperature of 740° C. for the duration of 2 hours and subsequently cooled down in an accelerated manner in water. This brought about the assimilation of the properties of the cast state of A and R with regard to strength and hardness.

(21) The degrees of cold forming s of 57% and 91% that are achievable for working example A underline the fact that the alloy of the invention, in spite of the content of hard particles, can achieve and even surpass the shape-changing properties of the conventional copper-nickel-tin alloy R.

(22) The thermal sensitivity of the reference material R with regard to the formation of the Sn-rich segregations was also found in the annealing between the two cold forming steps (No. 4 in Table 3). For this reason, the annealing temperature of 740° C. that was used for the intermediate annealing of the cold-rolled plate of alloy A had to be lowered to 690° C. for R.

(23) TABLE-US-00003 TABLE 3 Manufacturing program 1 for strips made from the strand-cast plates of working example A and of reference material R No. Manufacturing steps 1 Strand casting of plates of alloys A and R 2 Annealing the cast plate of alloy A: 740° C./2 h + water quench 3 Cold rolling Alloy A: from 11 to 4.70 mm (ε = 57%, φ = 0.8) Alloy R: from 24.5 to 12.1 mm (ε = 50%, φ = 0.7) 4 Annealing Alloy A: 740° C./2 h + water quench Alloy R: 690° C./2 h + water quench 5 Cold rolling Alloy A: from 4.70 to 0.4 mm (ε = 91%, φ = 2.4) Alloy R: from 12.1 to 2.33 mm (ε = 81%, φ = 1.6) 6 Age hardening: 300° C./4 h, 400° C./3 h, 450° C./3 h + air cooling

(24) After the performance of the manufacturing program 1, the indices of the strips of materials A and R were ascertained after the last cold rolling operation and on completion of the age hardening that are listed in Table 4.

(25) It becomes clear that the strengths and hardnesses of the strips of the working example A that have been cold-rolled and age-hardened at 300° C. are higher than the respective properties of the strips of the reference material R.

(26) Favored by the high content of hard particles, over and above the temperature of about 400° C., recrystallization of the microstructure of alloy A takes place. This recrystallization leads to a drop in strength and in hardness, and so the effect of the precipitation hardening and spinodal segregation cannot be manifested. Since no recrystallization of the microstructure is observed for the reference material R up to 450° C., the values for R.sub.m, R.sub.p0.2 and for the hardness, particularly after age hardening at 400° C., are higher for the reference material R than for the working example A.

(27) The microstructure of the further-processed working example A, after age hardening at 450° C., includes the hard particles of the second class (labeled 3 in FIG. 3).

(28) In addition, further phases have precipitated out in the microstructure of the further-processed alloy A. These include the continuous precipitates of the (Cu, Ni)—Sn system that are labeled 4 in FIG. 3, and the hard particles of the third class.

(29) The size of the hard particles of the third class of less than 3 μm is characteristic of the further-processed alloy of the invention. For the further-processed working example A of the invention, after age hardening at 450° C., it is actually less than 1 μm (labeled 5 in FIG. 4).

(30) TABLE-US-00004 TABLE 4 Grain size, electrical conductivity and mechanical indices of the cold-rolled and age-hardened strips of alloys A and R after undergoing manufacturing program 1 (Table 3) Age Electrical harden- Grain conducti- Hard- ing size vity R.sub.m R.sub.p0.2 A E ness Alloy [° C./h] [μm] [% IACS] [MPa] [MPa] [%] [GPa] HV1 A — — 10.6 974 917 3.2 119 311 300° C./4 h — 15.6 969 927 5.3 125 320 400° C./3 h .square-solid.<2 23.2 704 679 17.5 126 237 450° C./3 h  <1 24.5 591 575 22.3 126 193 R — — 10.7 838 787  7.2 120 267 300° C./4 h — 13.8 910 874  9.2 118 297 400° C./3 h — 22.0 793 735 13.6 108 264 450° C./3 h — 23.2 610 508 23.0 124 195 .square-solid. = not yet fully recrystallized

(31) In order to reduce the effect of the cold formability and the recrystallization temperature on the properties of the individual alloys, a further manufacturing program was conducted. This manufacturing program 2 pursued the aim of processing the strand-cast plates of materials A and R by means of cold-forming and annealing operations to give strips, using identical parameters in each case for the degrees of cold forming and the annealing temperatures (Table 5).

(32) Due to the high hardness of the cast state of the working example A, it was again calcined before the first cold rolling step at the temperature of 740° C. for the duration of 2 hours and subsequently cooled in an accelerated manner in water. As in the manufacturing program 1, this assimilated the properties of the cast state of A and R with regard to strength and hardness.

(33) TABLE-US-00005 TABLE 5 Manufacturing program 2 for strips made from the strand-cast plates of working example A and reference material R No. Manufacturing steps 1 Strand casting of plates of alloys A and R 2 Annealing of the cast plate of alloy A: 740° C./2 h + water quench 3 Cold rolling: from 9 to 6 mm (ε = 33%, φ = 0.4) 4 Annealing: 690° C./2 h + water quench 5 Cold rolling: from 6 to 3.5 mm (ε = 42%, φ = 0.5) 6 Annealing: 690° C./1 h + water quench 7 Cold rolling: from 3.5 to 3.0 mm (ε = 14%, φ = 0.15) 8 Age hardening: 400° C./3 h, 450° C./3 h, 500° C./3 h + air cooling

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

(35) The age hardening operation at 400° C. for three hours, due to the spinodal segregation of the microstructure, the rise in the strengths R.sub.m (from 498 to 717 MPa) and R.sub.p0.2 (from 439 to 649 MPa) and in the hardness HB (from 166 to 230 MPa) was at its clearest for the alloy R. However, the microstructure of the age-hardened states of the alloy R is very inhomogeneous with a grain size between 5 and 30 μm. Moreover, the microstructure of the age-hardened states of the reference material R is marked by discontinuous precipitates of the (Cu, Ni)—Sn system (labeled 1 in FIG. 1 and FIG. 2). Also present in the microstructure of the further-processed state of the reference material R are Ni phosphides (labeled 2 in FIG. 1 and FIG. 2).

(36) By contrast, the microstructure of the age-hardened strips of the working example A of the invention is very uniform with a grain size of 2 to 8 μm. Moreover, the structure of the working example A lacks the discontinuous precipitates even after age hardening at 450° C. for three hours followed by air cooling. By contrast, the hard particles of the second class are detectable in the microstructure. These phases are labeled 3 in FIG. 5 and FIG. 6.

(37) In addition, further phases have precipitated out in the microstructure of the further-processed alloy A. These include the continuous precipitates of the (Cu, Ni)—Sn system labeled 4 in FIG. 5 and the hard particles of the third class. For the further-processed working example A of the invention, the size of the hard particles of the third class after age hardening at 450° C. is even less than 1 μm (labeled 5 in FIG. 6).

(38) The strengths R.sub.m and R.sub.p0.2 of the strips of the alloy A after age hardening at 400° C./3 h/air, due to the spinodal segregation of the microstructure, assume the values of 690 and 618 MPa. Thus, R.sub.m and R.sub.p0.2 are lower than the indices of the correspondingly age-hardened state of the alloy R. The reason for this is that the working example A lacks the Ni content bound within the hard particles for the strength-increasing spinodal segregation of the microstructure. Should the strength level of the alloy R be a particular requirement, it is possible to add a higher proportion of the alloy element nickel to the alloy of the invention.

(39) TABLE-US-00006 TABLE 6 Grain size, electrical conductivity and mechanical indices of the cold-rolled and age-hardened strips of alloys A and R after undergoing manufacturing program 2 (Table 5) Age Electrical Hard- harden- Grain conducti- ness ing size vity R.sub.m R.sub.p0.2 A E HBW Alloy [° C./h] [μm] [% IACS] [MPa] [MPa] [%] [GPa] 1/30 A — — 11.6 556 498 25.1 113 188 400° C./3 h 2-8 15.1 690 618 21.4 132 222 450° C./3 h 2-8 16.8 666 534 22.1 126 211 500° C./3 h 2-8 16.7 614 444 24.4 124 190 R — — 11.2 498 439 27.9 104 166 400° C./3 h .square-solid.5-30 15.2 717 649 17.8 132 230 450° C./3 h .square-solid.5-30 17.0 705 591 20.6 121 219 500° C./3 h .square-solid.5-20 18.6 628 420 24.6 118 190 .square-solid. = inhomogeneous

(40) The next step included the testing of the hot formability of the strand casting of the alloys A and R. For this purpose, the cast plates were hot-rolled at the temperature of 720° C. (Table 7). For the further processing steps of cold forming and intermediate annealing, the parameters of manufacturing program 2 were adopted.

(41) Manufacturing program 3 for strips made from the strand-cast plates of the working example A and of the reference material R.

(42) TABLE-US-00007 TABLE 7 Manufacturing program 3 for strips made from the strand-cast plates of working example A and of reference material R No. Manufacturing steps 1 Strand casting of plates of alloys A and R 2 Alloy A, R: hot rolling at 720° C. + water quench 3 Cold rolling of alloy A: from 9 to 6 mm (ε = 33%, φ = 0.4) 4 Annealing of alloy A: 690° C./2 h + water quench 5 Cold rolling of alloy A: from 6 to 3.5 mm (ε = 42%, φ = 0.5) 6 Annealing of alloy A: 690° C./1 h + water quench 7 Cold rolling of alloy A: from 3.5 to 3.0 mm (ε = 14%, φ = 0.15) 8 Age hardening of alloy A: 400° C./3 h, 450° C./3 h + air cooling

(43) During the hot rolling of the cast plates of the reference alloy R, deep heat cracks formed even after a few passes, which led to failure of the plates through fracture.

(44) By contrast, the cast plates of the working example A of the invention were hot-rollable without damage and could be manufactured to the final thickness of 3.0 mm after multiple cold rolling and calcination processes. The properties of the age-hardened strips (Table 8) correspond largely to those of the strips that have been produced without hot forming by the manufacturing program 2 (Table 6).

(45) Also comparable is the microstructure of the strips made from the working example A of the alloy of the invention that were manufactured without and with a hot forming step. Thus, FIG. 7 and FIG. 8 show the uniform structure of the strips made from the working example A that were produced with a hot forming stage and a subsequent age hardening operation at 400° C./3 h/air cooling. In FIG. 7 and FIG. 8, the hard particles of the second class, labeled 3, are again apparent.

(46) In addition, FIG. 7 shows the continuous precipitates of the (Cu, Ni)—Sn system, labeled 4, and the hard particles of the third class. In the microstructure of the further-processed variant of the working example A, the hard particles of the third class actually assume a size of less than 1 μm (labeled 5 in FIG. 8).

(47) The analysis of the hard particles of the second and third class in this further-processed state of the working example A revealed the compound SiB.sub.6 as a representative of the Si-containing and B-containing phases, Ni.sub.6Si.sub.2B as a representative of the Ni—Si borides, Ni.sub.3B as a representative of the Ni borides, FeB as a representative of the Fe borides, Ni.sub.3P as a representative of the Ni phosphides, Fe.sub.2P as a representative of the Fe phosphides, Ni.sub.2Si as a representative of the Ni silicides, and Fe-rich particles, which are present individually and as addition compounds and/or mixed compounds in the microstructure. In addition, these hard particles are ensheathed by precipitates of the (Cu, Ni)—Sn system.

(48) TABLE-US-00008 TABLE 8 Grain size, electrical conductivity and mechanical indices of the cold-rolled and age-hardened strips of working example A after undergoing manufacturing program 3 (Table 7) Age Electrical Hard- harden- Grain conducti- ness ing size vity R.sub.m R.sub.p0.2 A E HBW Alloy [° C./h] [μm] [% IACS] [MPa] [MPa] [%] [GPa] 1/30 A — — 1.8 554 501 23.8 110 185 400° C./3 h 3-10 15.3 679 610 21.8 127 217 450° C./3 h 3-10 16.8 658 535 20.8 126 205

(49) The subsequent test stage comprised the testing of the hot forming characteristics of the working example 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 the manufacturing program 3. This measure enabled the study of the cold formability of the hot-rolled strip state of the alloy A. The individual process steps of the manufacturing program 4 are apparent from Table 9.

(50) TABLE-US-00009 TABLE 9 Manufacturing program 4 for strips from the strand-cast plates of working example A No. Manufacturing steps 1 Strand-casting of plates of alloy A 2 Alloy A: hot rolling at 780° C. + water quench 3 Cold rolling of alloy A: from 9 to 1.4 mm (ε = 84%, φ = 1.9) 4 Annealing of alloy A: 690° C./1 h + water quench 5 Cold rolling of alloy A: from 1.4 to 1.2 mm (ε = 14%, φ = 0.15) 6 Age hardening: 350° C./3 h, 400° C./3 h, 450° C./3 h, 500° C./ 3 h + air cooling

(51) At the higher hot rolling temperature, the strand-cast plates of the alloy A showed excellent hot formability. The hot-rolled plates were also cold-rollable without difficulty with an extremely high degree of cold forming s of 84%. In order to be able to make the age hardening outcome comparable with the result of the preceding manufacturing program 3, the last cold rolling step followed after recrystallization annealing at 690° C. with the same degree of cold forming ε of 14%.

(52) After the strips had been age-hardened within the temperature range from 350 to 500° C., the grain size of the very uniform microstructure was 5 to 10 μm (Table 10). Particularly at the age hardening temperature of 400° C., the spinodal segregation of the microstructure of the alloy of the invention leads to a marked rise in strength and hardness. For instance, there is a rise in the tensile strength R.sub.m from 557 MPa in the cold-rolled state to 692 MPa in the age-hardened state. There is also a rise in the hardness HB from 177 to 210.

(53) TABLE-US-00010 TABLE 10 Grain size, electrical conductivity and mechanical indices of the cold-rolled and age-hardened strips of alloy A after undergoing manufacturing program 4 (Table 9) Age Electrical Hard- harden- Grain conducti- ness ing size vity R.sub.m R.sub.p0.2 A E HBW Alloy [° C./h] [μm] [% IACS] [MPa] [MPa] [%] [GPa] 1/30 A — — 11.6 557 520 22.2 — 177 350° C./3 h 5-10 13.8 674 607 23.3 143 204 400° C./3 h 5-10 15.2 692 614 20.1 150 210 450° C./3 h 5-10 17.2 659 519 21.9 128 193 500° C./3 h 5-10 15.8 598 437 25.0 128 170

(54) In the construction of installations, devices, engines and machinery, components having relatively high dimensions are required for numerous applications. For example, this is often the case in the field of slide bearings. The production of the corresponding components requires a precursor material of appropriately large dimensions. Therefore, due to the limited producibility of infinitely large castings, it is necessary to establish the required material properties if at all possible by means of small degrees of cold forming as well.

(55) Table 11 lists the process steps that are used in the course of the manufacturing program 5. The manufacturing operation was carried out with one cycle of cold forming and annealing operations. Again, only the cast plates of the alloy A were calcined prior to the first cold rolling operation at 740° C.

(56) The first cold rolling operation on the cast plate of the alloy R and on the annealed cast plate of the alloy A was implemented with a degree of forming s of 16%. An annealing operation at 690° C. was followed by a cold rolling operation with e of 12%. Finally, age hardening of the strips took place at the temperatures of 350° C., 400° C. and 450° C.

(57) TABLE-US-00011 TABLE 11 Manufacturing program 5 for strips from the strand-cast plates of working example A and of reference material R No. Manufacturing steps 1 Strand casting of plates of alloys A and R 2 Annealing of the cast plates of alloy A: 740° C./2 h + water quench 3 Cold rolling alloy A, R: from 9 to 7.6 mm (ε = 16%, φ = 0.17) 4 Annealing alloy A, R: 690° C./2 h + water quench 5 Cold rolling alloy A, R: from 7.6 to 6.7 mm (ε = 12%, φ = 0.126) 6 Age hardening: 350° C./3 h, 400° C./3 h, 450° C./ 3 h + air cooling

(58) The low degree of cold forming in the first cold rolling step of ε=16%, together with the subsequent annealing operation at 690° C., was insufficient to eliminate the dendritic and coarse-grain microstructure of the reference material R. Moreover, this thermomechanical treatment enhanced the coverage of the grain boundaries of the alloy R with Sn-rich segregations.

(59) Across the dendritic structure and across the grain boundaries of the alloy R covered by Sn-rich segregations, cracks running from the surface deep into the interior of the strip formed during the second cold rolling step.

(60) The crack-free and homogeneous microstructure of the strips of the working example A is characterized by the arrangement of the hard particles of the second and third class. As was already the case after the preceding manufacturing programs, the hard particles of the third class have a size of less than 1 μm, even after this manufacturing program 5.

(61) The resulting properties of the strips after the last cold rolling operation and after the age hardening operation are shown in Table 12. Due to the high density of cracks, it was not possible to take undamaged tensile samples from the strips of the material R. Thus, it was possible to undertake only the metallographic analysis and the measurement of hardness on these strips.

(62) The working example A has a high degree of age hardenability which is manifested by the interaction of the mechanisms of precipitation hardening and spinodal segregation of the microstructure. Thus, there is a rise in the indices R.sub.m and R.sub.p0.2 as a result of age hardening at 400° C. from 518 to 633 MPa and from 451 to 575 MPa.

(63) TABLE-US-00012 TABLE 12 Grain size, electrical conductivity and mechanical indices of the cold-rolled and age-hardened strips of alloys A and R after undergoing manufacturing program 5 (Table 11) Electrical Hard- Age Grain conducti- ness harden- size vity R.sub.m R.sub.p0.2 A E HBW Alloy ing [μm] [% IACS] [MPa] [MPa] [%] [GPa] 1/30 A — — 11.6 518 451 21.8 124 188 350° C./3 h 20 13.4 610 530 23.5 130 212 400° C./3 h 20-25 14.4 633 575 19.1 116 217 450° C./3 h 20-25 16.0 630 496 17.4 108 204 R — .square-solid.— Not possible owing to formation of 175 350° C./3 h .square-solid.— cracks! 242 400° C./3 h .square-solid.— 229 450° C./3 h .square-solid.— 217 .square-solid. = dendritic, with Sn-rich segregations

(64) As a result, it can be stated that, by means of the variation of the chemical composition, the degrees of forming for the cold forming operation(s) and the variation in the age hardening conditions, it is possible to adjust the degree of precipitation hardening and the degree of spinodal segregation of the microstructure of the invention to the required material properties. In this way, it is possible to bring the strength, hardness, ductility and electrical conductivity of the alloy of the invention into line with the field of use envisaged.