T7 Solution treated, quenched, and overaged/stabilized Wrought products that are overaged to carry them beyond a point of maximum strength to provide control of some other characteristic, usually corrosion resistance. Applies to cast products that are artificially aged after solution heat treatment to provide dimensional and strength stability. T8 Solution treated, quenched, cold worked, and artificially aged Products that are cold worked after solution treatment and the effect of cold work is recognized in mechanical property limits. T9 Solution treated, quenched, artificially aged, and then cold worked Products that are cold worked to improve strength. Table 7 Typical tensile properties of selected 2xxx, 6xxx, and 7xxx aluminum alloy products Tensile strength Yield strength Alloy Temper Product (a) MPa ksi MPa ksi Elongation in 50 mm, % T4 Sheet 250 36 125 18 28 2008 T6 Sheet 300 44 240 35 13 2014 T6, T651 Plate, forging 485 70 415 60 13 T3, T351 Sheet, plate 450 65 310 45 18 T361 Sheet, plate 495 72 395 57 13 T81, T851 Sheet, plate 485 70 450 65 6 2024 T861 Sheet, plate 515 75 490 71 6 2224 T3511 Extrusion 530 77 400 58 16 2324 T39 Plate 505 73 415 60 12 2524 T3, T351 Sheet, plate 450 65 310 45 21 2036 T4 Sheet 340 49 195 28 24 T81, T851 Sheet, plate 455 66 350 51 10 2219 T87 Sheet, plate 475 69 395 57 10 2519 T87 Plate 490 71 430 62 10 T4 Sheet 220 32 125 18 25 6009 T62 Sheet 300 44 260 38 11 T4 Sheet 285 41 165 24 25 6111 T6 Sheet 350 51 310 45 10 T6, T6511 Sheet, plate, extrusion, forging 310 45 275 40 12 6061 T9 Extruded rod 405 59 395 57 12 T5 Extrusion 185 27 145 21 12 6063 T6 Extrusion 240 35 215 31 12 7005 T5 Extrusion 350 51 290 42 13 7049 T73 Forging 540 78 475 69 10 7050 T74, T745X Plate, forging, extrusion 510 74 450 65 13 T651, T6151 Plate 600 87 560 81 11 7150 T77511 Extrusion 650 94 615 89 12 T7751 Plate 640 93 615 89 10 7055 T77511 Extrusion 670 97 655 95 11 T6, T651 Sheet, plate 570 83 505 73 11 7075 T73, T735X Plate, forging 505 73 435 63 13 T7351 Plate 505 73 435 63 15 7475 T7651 Plate 455 66 390 57 15 (a) Properties of sheet and plate are for the long-transverse direction, and those of extrusions and forgings are for the longitudinal direction. Table 8 Typical tensile properties of selected aluminum cast products used in the automotive industry Tensile strength Yield strength Alloy Temper Casting process MPa ksi MPa ksi Elongation in 20-25 mm, % 354.0 T6 Permanent mold 372 54 255 57 8 Sand 276 40 207 30 6 A356.0 T6 Permanent mold 283 41 207 30 12 365.0 T6 Vacuum die 206 30 139 20 17 The mechanical properties of a casting can vary due to composition, geometry, and fabric ation process, which are selected based on design requirements. Properties in certain areas can be adjusted by process and tooling design. Table 9 Typical tensile properties of selected 3xxx and 5xxx aluminum alloy sheet products Tensile strength Yield strength Alloy Temper MPa ksi MPa ksi Elongation in 50 mm, % 0 100 16 40 6 30 H14 125 18 115 17 9 3003 H18 165 24 150 22 5 O 180 26 70 10 20 H34 240 35 200 29 9 H38 285 41 250 36 5 3004 H19 295 43 285 41 2 O 130 19 55 8 25 3005 H14 180 26 165 24 7 H18 240 35 225 33 4 O 115 17 55 8 24 H25 180 26 160 23 8 3105 H18 215 31 195 28 3 O 125 18 40 6 25 H34 160 23 140 20 8 5005 H38 200 29 185 27 5 O 145 21 55 8 24 H34 190 28 165 24 8 5050 H38 220 32 200 29 6 O 195 28 90 13 25 H34 260 38 215 31 10 5052 H38 290 42 255 37 7 O 180 26 85 12 23 H25 235 34 170 25 11 5252 H28 285 41 240 35 5 O 240 35 115 17 27 H34 290 42 230 33 13 H38 330 48 270 39 10 5154 H112 240 35 115 17 25 5454 O 250 36 115 17 22 H34 305 44 240 35 10 H111 250 36 125 18 18 H112 250 36 125 18 18 O 290 42 150 22 35 H18 435 63 405 59 10 5056 H38 415 60 345 50 15 O 310 45 160 23 24 H112 310 45 165 24 22 5456 H116 350 51 255 37 16 5182 O 275 40 130 19 21 O Temper. Ductility of all aluminum products is highest in this temper. Factors determining properties of annealed products depend on the alloy system. The annealed strength of unalloyed aluminum, 1xxx series, generally increases and ductility decreases with increasing impurity level, and the amounts of magnesium and manganese largely determine the bulk properties of 3xxx and 5xxx products. Strength increases as magnesium (solid-solution strengthening) and manganese contents (dispersion and solid-solution strengthening) increase. Ultimate tensile strength increases more significantly with increasing magnesium content than does tensile yield strength because of the potent effect of magnesium on work hardening. Parts of heat-treatable alloy products that are difficult to form are often formed in the O temper, then heat treated to a final temper. H Tempers. Cold working of annealed material to H1 tempers increases the dislocation density. This increases strength, particularly yield strength, and decreases ductility. In unalloyed aluminum and in alloys containing little magnesium, cold working produces cells that have walls containing a high density of dislocations enclosing a volume of relatively strain- free material. In alloys containing sufficient amounts of magnesium, however, the dislocations form a tangled forest. In highly worked aluminum-magnesium alloys, rearrangement of the dislocation structure occurs over long times at room temperature. Stabilizing treatments, H3 tempers, prevent loss of strength in certain 3xxx and 5xxx alloys during subsequent long-time exposure. During partial annealing treatments, H2 tempers, cell walls either form or become more perfect, and dislocations within the cell migrate to cell boundaries. If the temperature exceeds a critical level, which depends on alloy content and strain, the cold-worked product will either partially or completely recrystallize. Materials in H2 tempers provide a combination of strength and ductility generally superior to that of material in H1 tempers. Cold- worked alloys containing above approximately 3.5% Mg and annealed alloys containing above approximately 4.5% Mg can also suffer a degradation in corrosion resistance caused by precipitation of a continuous film of Al 3 Mg 2 on grain boundaries at temperatures between ambient and approximately 205 °C (400 °F). Special H116 and H321 controlled hot- rolling tempers have been developed that either avoid precipitation of Al 3 Mg 2 on grain boundaries or agglomerate the precipitate to increase corrosion resistance. For a particular strength level, a higher resistance to stress corrosion is obtained by increasing magnesium and manganese rather than by increasing work hardening. W and T Tempers. The highest strengths are obtained by precipitation hardening. The material is held for a sufficient time above the solvus to dissolve essentially all of the major alloying elements, quenched at a rate to retain most of these elements in solid solution, then aged either at room temperature (natural aging) or at a modestly elevated temperature (artificial aging). The highest-strength alloys contain the largest concentration of the major alloying elements. For a particular alloy system, strength typically increases with increasing alloy content. Most 2xxx and 6xxx wrought alloys and 2xx.x and 3xx.x cast alloys are strengthened during natural aging by Guinier-Preston (G-P) zones, which are precursors to Al 2 Cu, Al 2 CuMg, Mg 2 Si, or Al 4 CuMg 5 Si 4 phases. Strength of these materials increases for about 4 days, then stabilizes (T4 temper). In contrast, in 7xxx alloys containing G-P zone precursors to phases such as MgZn 2 , strength continues to increase indefinitely at room temperature (W temper). The ductility in the freshly quenched (W < h temper) condition is high enough for many forming operations. Consequently, many parts are formed shortly after quenching from the solution-heat-treatment temperature. To prevent the formation of large grains during the solution treatment of formed parts, a critical amount of strain must be avoided. Although this critical strain is alloy dependent, strains near 10% are particularly troublesome for most alloys. In addition, ductility in the T4 or T3 tempers is sufficiently high that some parts can be formed successfully in this condition. Strength, particularly yield strength, increases substantially with artificial aging (T6 temper). This increase is accompanied by a loss in ductility. Strength of materials hardened by Al 2 CuMg, Al 2 Cu, or Al 2 CuLi precipitates may be increased by cold work prior to artificial aging, T8, treatments. The increase in strength of these materials is attributed to a refinement of Al 2 CuLi and of the metastable precursors to Al 2 CuMg and Al 2 Cu. Additions of silicon and other alloying elements can also serve to refine the size of precipitates in certain 2xxx alloys. Cold-finishing rod and bar products after artificial aging increases their strength (T9 temper). The solution treatment, in most cases, is a separate operation. In particular circumstances, however, the heat from a shaping process may be sufficient to provide solution treatment. These products can be cooled after the shaping process and subsequently aged to develop useful properties (T5 temper). Some 6xxx alloys attain the same specified properties whether furnace solution heat treated or cooled from an elevated-temperature shaping process at a rate rapid enough to maintain sufficient silicon and magnesium in solution. In such cases, the T6 temper designation may be used. Aluminum Alloy Microstructural Features Not Inferred from the Alloy-Temper Designation Systems The alloy designation system defines the alloy content, and the temper-designation system identifies many of the thermal and mechanical processes that control the microstructure and, hence, the bulk properties of aluminum alloy products. Nevertheless, many metallurgical features are not specified by these systems. The features include nonmetallic inclusions, porosity, second-phase particles, grain and dislocation structure, and crystallographic texture. Inclusions are typically oxides of aluminum and magnesium including spinel, MgAl 2 O 4 . Oxides form on the surface of molten aluminum and become entrapped when turbulent flow forces them below the surface. Filtration of the molten metal is used to control inclusions. Inclusions can give rise to problems ranging from pinholes in foil to reduced fatigue life in structural wrought products and castings. Porosity reduces ductility and increases susceptibility to the initiation of fatigue cracks. Porosity may arise from either shrinkage during solidification or from hydrogen. Hydrogen control during solidification is extremely important because of the ten-fold decrease in the solubility of hydrogen in aluminum as it solidifies. Hydrogen-induced porosity can also occur in solid aluminum products when they are heated to high temperatures in humid environments. Provided that the hydrogen content is low enough, most of the porosity can be closed by thermomechanical treatments. Isostatic pressure can be used to close the pores in castings, and conventional forging and extrusion are effective in healing ingot porosity. Porosity in thick-rolled products is particularly difficult to close; tensile stresses in the short-transverse direction may arise during the initial rolling of thick plate because the amount of deformation per pass is limited. This stress causes pores to enlarge. With additional rolling to thinner plate, the pores heal. Second-phase particles are divided into four classes based on their mode of formation and their ability to be dissolved: primary particles, constituents, dispersoids, and precipitates. Primary Particles. These particles form when some phase other than aluminum solid solution separates first from the melt. Primary silicon particles form in castings when hypereutectic aluminum-silicon alloys solidify by eutectic decomposition. Ductility decreases with increasing size of the silicon particles, so size control is important. The coarse, faceted primary silicon particles are refined to a fine spherulitic structure using additives containing phosphorus. In certain casting alloys and 8xxx wrought alloys, primary iron-bearing constituents can form if the alloying content is such that the alloy is hypereutectic. In wrought alloys, macroscopically large, undesirable primary particles of Al 7 Cr, Al 3 Ti, or Al 3 Zr can form by a peritectic reaction if chemical composition is not closely controlled. Constituents. These particles may be either intermetallic compounds or essentially pure silicon that forms during solidification of hypoeutectic aluminum-silicon alloys. They range in size from a few micrometers to tens of micrometers. Constituents can be classified either as virtually insoluble or soluble. Because the low maximum solid solubility of iron in aluminum is further reduced by other alloying elements to 0.01 wt% or less, constituents containing iron are insoluble. Iron-free constituents containing silicon can be either soluble or insoluble depending on the chemical composition of the alloy. Major alloying elements can combine either with each other or with aluminum to form soluble constituent particles. Most of these soluble constituents dissolve either during ingot preheating prior to deformation processing or during the solution heat treatment of cast shapes or wrought products. Constituent size decreases with increasing solidification rate. In hypoeutectic 3xx.0 and 4xx.0 castings, modification by elements such as strontium significantly refine the flake structure of the silicon particles to a finer fibrous morphology. Constituent particles are generally not beneficial and are detrimental to the fatigue resistance and fracture toughness of high-strength alloy products. These particles fracture at relatively low plastic strains and provide low-energy sites for the initiation and growth of cracks. Several high-purity (low iron and silicon) versions of 2024 and 7075 have been commercialized, and the maximum allowable impurity levels of all modern high-strength alloys are significantly lower than those of older alloys. Despite the harmful effects of constituents in high-strength alloys, the ability of alloy 3004- H19 to make commercially successful beverage containers relies on careful control of size, volume fraction, and distribution of Al 12 (Fe,Mn)Si constituent particles. These constituent particles serve to "scour" the die during the drawing operation so that galling is minimized. Attempts to produce can stock from roll-cast sheet have generally not been successful because the particle size distribution in roll-cast sheet is not as effective in minimizing galling. Dispersoids. These form by solid-state precipitation, either during ingot preheating or during the thermal heat treatment of cast shapes, of slow-diffusing supersaturated elements that are soluble in molten aluminum but which have limited solubility in solid aluminum. Manganese, chromium, or zirconium are typical dispersoid-forming elements. Unlike the precipitates that confer precipitation hardening, dispersoids are virtually impossible to dissolve completely, once precipitated. In addition to providing dispersion strengthening, the size distribution of dispersoids in wrought alloys are a key factor in controlling degree of recrystallization, recrystallized grain size, and crystallographic texture. Dispersoids in non-heat-treatable alloys also stabilize the deformation substructure during elevated-temperature exposures, for example, during paint baking. In contrast to the commercially significant dispersion strengthening provided by dispersoids in 3xxx and 5xxx alloys, the level of dispersion strengthening afforded by dispersoids in wrought heat-treatable alloys is trivial. In 2x24 alloys, Al 20 Cu 2 Mn 3 dispersoids nucleate dislocations at the particle-matrix interface during the quench. These dislocations serve as nucleation sites for subsequent precipitation. The newer 7xxx alloys contain zirconium, which forms coherent Al 3 Zr dispersoids while most of the older 7xxx alloys contain Al 12 Mg 2 Cr dispersoids which exhibit incoherent interfaces. The incoherent interfaces serve to nucleate MgZn 2 precipitates during the quench, so alloys containing these precipitates lose a great deal of their potential to develop high strength after slow quenching (quench sensitivity). Nucleation is difficult on coherent interfaces, so the newer alloys are less quench sensitive. A number of casting alloys, and some wrought alloys, contain elements that can form either constituents or dispersoids depending on the solidification rate. Precipitates can form during any thermal operation below the solvus. In properly solution-heat-treated products, all precipitates dissolve during the solution-heat-treatment operation. Depending on quench rate and alloy, precipitates can form during the quench from the solution-heat-treatment temperature at grain and subgrain boundaries and at particle- matrix interfaces. These coarse precipitates do not contribute to age hardening and can serve to reduce properties such as ductility, fracture toughness, and resistance to intergranular corrosion. After the quench, G-P zones form at ambient temperature (natural aging). These are agglomerates of atoms of the major solute elements with a diffuse, coherent boundary between the G-P zone and the matrix. During elevated-temperature precipitation heat treatments (artificial aging) G-P zones may either nucleate metastable precipitates or they may dissolve, and metastable precipitates nucleate separately. Cold working subsequent to quenching introduces dislocations that may serve to nucleate metastable or equilibrium precipitates. With prolonged artificial aging, equilibrium precipitates may form. Coarse equilibrium precipitates form during annealing treatments of heat-treatable alloy products, O temper. They also form during most thermomechanical treatments prior to solution heat treatment. Grain Structure. The grain size of aluminum alloy ingots and castings is typically controlled by the introduction of inoculants that form intermetallic compounds containing titanium and/or boron. During deformation processing, the grain structure becomes modified. Most aluminum alloy products undergo dynamic recovery during hot working as the dislocations form networks of subgrains. New dislocation-free grains may form between and following rolling passes (static recrystallization) or during deformation processing (dynamic recrystallization). During deformation, the crystal lattice of the aluminum matrix rotates at its interfaces between constituent and coarse precipitate particles. These high- energy sites serve to nucleate recrystallization. This process is termed particle-stimulated nucleation and is an important mechanism in the recrystallization process of aluminum. The particle size that will serve as a nucleus decreases as deformation temperature decreases and strain and strain rate increase. Dispersoid particles retard the movement of high- angle grain boundaries. Consequently, hot-worked structures are resistant to recrystallization and often retain the dynamically recovered subgrain structure in the interiors of elongated cast grain boundaries. In heat-treated products containing a sufficient quantity of dispersoids the unrecrystallized structure of hot-worked-plate forgings and extrusions can be retained after solution heat treatment. Degree of recrystallization of hot-worked products has an effect on fracture toughness. Unrecrystallized products develop higher toughness than do products that are either partially or completely recrystallized. This behavior is attributed to precipitation on the recrystallized high-angle grain boundaries during the quench. These particles increase the tendency for low-energy intergranular fracture. Products such as sheet, rods, and tubing that are cold rolled invariably recrystallize during solution heat treatment or annealing to O temper. Decreasing the grain size can increase strength of 5xxx alloy products in the O temper by 7 to 28 MPa (1 to 4 ksi), but grain size is not a major factor in increasing strength of other aluminum alloy products. Several measures of formability are influenced by grain size, however, so grain size is controlled for this reason. One particular use of grain size control is to produce stable, fine grains, which are essential in developing superplastic behavior in aluminum alloy sheet. Crystallographic Texture. Cast aluminum ingots and shapes generally have a random crystallographic texture; the orientation of the unit cells comprising each grain are not aligned. With deformation, however, certain preferred crystallographic orientations develop. Many of the grains rotate and assume certain orientations with respect to the direction of deformation. For flat-rolled products and extrusions having a high aspect ratio of width to thickness, the deformation texture is similar to that in pure fcc metals. These orientations are described by using the Miller indices of the planes {nnn} in the grains parallel to the plane of the worked product and directions [nnn] parallel to the working direction. The predominant textures are {110}[112], {123}[634], and {112}[111]. During recrystallization, a high concentration of grains in the {001}[100] or {011}[100] orientations may develop. Alternatively, if particle-stimulated nucleation is present to a large extent, the recrystallized texture will be random. Control of crystallographic texture is particularly important for non-heat-treatable sheet that will be drawn. If texture is not random, ears form during the drawing process. In extruded or drawn rod or bar, the texture is a dual-fiber texture in which almost all grains are aligned so that the grain directions are either [001] or [111]. In heat-treatable alloys, texture has the most potent effect on the properties of extrusions that have the dual-fiber texture. Strengthening by this process is so potent that the longitudinal yield strengths of extruded products exhibiting this texture are about 70 MPa (10 ksi) higher than strength in the transverse direction. If this dual-fiber texture is lost by recrystallization, strength in the longitudinal direction decreases to that in the transverse directions. References cited in this section 2. H. Baker, Ed., Alloy Phase Diagrams, Vol 3, ASM Handbook, ASM International, 1992 3. J.R. Davis, Ed., ASM Specialty Handbook: Aluminum and Aluminum Alloys, ASM International, 1993 4. D. Altenpohl, Aluminum Viewed from Within, an Introduction to the Metallurgy of Aluminum Fabrication, Aluminium-Verlag, Dusseldorf, 1982 5. Aluminum Standards and Data, The Aluminum Association, 1993 6. C. Brooks, Heat Treatment, Structure and Properties of Nonferrous Alloys, American Society for Metals, 1982 7. J. Hatch, Ed., Aluminum: Properties and Physical Metallurgy, American Society for Metals, 1984 8. W.E. Haupin and J.T. Staley, Aluminum and Aluminum Alloys, Encyclopedia of Chemical Technology, 1992 9. Heat Treating of Aluminum Alloys, Heat Treating, Vol 4, ASM Handbook, ASM International, 1991, p 841-879 10. W. Petzow and G. Effenberg, Ed., Ternary Alloys: A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams, VCH Verlagsgesellschaft, Weinheim, Germany, 1990 11. H.W.L. Phillips, Equilibrium Diagrams of Aluminium Alloy Systems, Aluminum Development Association, 1961 12. I.J. Polmear, Light Alloys, Metallurgy of the Light Metals, 3rd ed., Arnold, 1995 13. R.E. Sanders, Jr., S.F. Baumann, and H. Stumpf, Non-Heat-Treatable Aluminum Alloys, Aluminum Alloys, Their Physical and Mechanical Properties, Engineering Materials Advisory Services Ltd, 1986, p 1441- 1484 14. T.H. Sanders, Jr., and J.T. Staley, Review of Fatigue and Fracture Research on High- Strength Aluminum Alloys, Fatigue and Microstructure, American Society for Metals, 1979, p 467-522 15. J.T. Staley, Metallurgical Factors Affecting Strength of High Strength Alloy Products, Pro ceedings of Fourth International Conference on Aluminum Alloys, Norwegian Institute of Technology, Department of Metallurgy and SINTEF Metallurgy, 1994 16. E.A. Starke, Jr., and J.T. Staley, Application of Modern Aluminum Alloys to Aircraft, Progr. Aerosp. Sci., Vol 32 (No. 2-3), 1996, p 131-172 17. K.R. Van Horn, Ed., Aluminum, Vol I, Properties, Physical Metallurgy and Phase Diagrams, American Society for Metals, 1967 Effects of Composition, Processing, and Structure on Properties of Nonferrous Alloys Ronald N. Caron, Olin Corporation; James T. Staley, Alcoa Technical Center Copper and Copper Alloys After iron and aluminum, copper is the third most-prominent commercial metal because of its availability and attractive properties: excellent malleability (or formability), good strength, excellent electrical and thermal conductivity, and superior corrosion resistance (Ref 18, 19, 20, 21, 22, 23, 24, 25). Copper offers the designer moderate levels of density (8.94 g/cm 3 , or 0.323 lb/in. 3 ), elastic modulus (115 GPa, or 17 × 10 6 psi), and melting temperature (1083 °C, or 1981 °F). It forms many useful alloys to provide a wide variety of engineering property combinations and is not unduly sensitive to most impurity elements. The electrical conductivity of commercially available pure copper, about 101% IACS (International Annealed Copper Standard), is second only to that of commercially pure silver (about 103% IACS). Standard commercial copper is available with higher purity and, therefore, higher conductivity than what was available when its electrical resistivity value at 20 °C (70 °F) was picked to define the 100% level on the IACS scale in 1913. The thermal conductivity for copper is also high, 391 W/m · K (226 Btu/ft · h · °F), being directly related to the electrical conductivity through the Wiedemann-Franz relationship. Copper and the majority of its alloys are highly workable hot or cold, making them readily commercially available in various wrought forms: forgings, bar, wire, tube, sheet, and foil. In 1995, copper used in wire and cable represented about 50% of U.S. production and in flat products of various thickness another 15%, rod and bar about 14%, tube about 14.5%, with foundries using about 5% for cast products, and metal powder manufacturers about 0.6%. Besides the more familiar copper wire, copper and its alloys are used in electrical and electronic connectors and components, heat-exchanger tubing, plumbing fixtures, hardware, bearings, and coinage. As with other metal systems, copper is intentionally alloyed to improve its strength without unduly degrading ductility or workability. However, it should be recognized that additions of alloying elements also degrade electrical and thermal conductivity by various amounts depending on the alloying element, its concentration and location in the microstructure (solid solution or dispersoid). The choice of alloy and condition is most often based on the trade-off between strength and conductivity. Alloying also changes the color from reddish brown to yellow (with zinc, as in brasses) and to metallic white or "silver" (with nickel, as in U.S. cupronickel coinage). Copper and its alloys are readily cast into cake, billet, rod, or plate suitable for subsequent hot or cold processing into plate, sheet, rod, wire, or tube via all the standard rolling, drawing, extrusion, forging, machining, and joining methods. Copper and copper alloy tubing can be made by the standard methods of piercing and tube drawing as well as by the continuous induction welding of strip. Copper is hot worked over the temperature range 750 to 875 °C (1400 to 1600 °F), annealed between cold working steps over the temperature range 375 to 650 °C (700 to 1200 °F), and is thermally stress relieved usually between 200 and 350 °C (390 and 660 °F). Copper and its alloys owe their excellent fabricability to the face-centered cubic crystal structure and its twelve available dislocation slip systems. Many of the applications of copper and its alloys take advantage of the work-hardening capability of the material, with the cold processing deformation of the final forming steps providing the required strength/ductility for direct use or for subsequent forming of stamped components. Copper is easily processible to more than 95% reduction in area. The amount of cold deformation between softening anneals is usually restricted to 90% maximum to avoid excessive crystallographic texturing, especially in rolling of sheet and strip. Although copper obeys the Hall-Petch relationship and grain size can be readily controlled by processing parameters, work hardening is the only strengthening mechanism used with pure copper. Whether applied by processing to shape and thickness, as a rolled strip or drawn wire, or by forming into the finish component, as an electrical connector, the amount of work hardening applied is limited by the amount of ductility required by the application. Worked copper can be recrystallized by annealing at temperatures as low as 250 °C (480 °F), depending on prior degree of cold work and time at temperature. While this facilitates processing, it also means that softening resistance during long-time exposures at moderately elevated temperatures can be a concern, especially in electrical and electronic applications where I 2 R heating is a factor. For applications above room temperature, but at temperatures lower than those inducing recrystallization in commercial heat treatments, thermal softening can occur over extended periods and characteristics such as the half- softening temperature should be considered; that is, the temperature for which the worked metal softens to half its original hardness after a specific exposure time, usually 1 h. A more useful engineering property for many electrical contact applications is stress-relaxation resistance, the property that characterizes the decrease in contact load supported by a mechanical contact over time at a given temperature, typically measured at room temperature between exposures to elevated temperature (Ref 20). Figure 3 illustrates the characteristics of the tensile-stress-relaxation property of drawn (worked) copper wire; the degree of relaxation increases with temperature and time. It also increases with the initial temper or degree of cold work in the material. The mechanism is the thermally activated and applied-stress directed motion of crystal lattice defects, such as point defects and dislocations. Consequently, the application of a thermal heat treatment (stabilization anneal) to induce recovery mechanisms to tie up mobile components of dislocations will improve the stress-relaxation resistance. Alloying elements also restrict dislocation motion and provide a more potent remedy for improving stress-relaxation resistance of cold- worked metal in service. For example, the improvement in stress-relaxation resistance obtainable by alloying copper with 5% Sn (alloy C51000) in combination with a low-temperature stabilization heat treatment and as a function of sheet orientation is illustrated by comparing the alloy data at 93 °C (200 °F) in Fig. 4 with those for copper in Fig. 3. Fig. 3 Tensile-stress- relaxation characteristics of copper alloy C11000. Data are for tinned 30 AWG (0.25 mm diam) annealed ETP copper wire; initial elastic stress, 89 MPa (13 ksi). [...]... Nickel-chromium-iron alloys Alloy 600 Ni-15Cr-8Fe 655 95 310 45 40 75 HRB Alloy 800 Ni-21Cr-39.5Fe-0.4Ti-0.4Al 600 87 295 43 44 138 HB Alloy 6 17 Ni-22Cr-3Fe-12Co-9Mo-1Al 75 5 110 350 51 58 173 HB Alloy 690 Ni-29Cr-9Fe 72 5 105 348 50 41 88 HRB Alloy 75 1 Ni-15Cr-7Fe-1Nb-2Ti 1310 190 976 142 22 352 HB Nickel-chromium-molybdenum alloys Alloy C- 276 Ni-15.5Cr-16Mo-5.5Fe-3 .75 W-1.25Co + V 78 5 114 372 54 62 209 HB Alloy... 301 Ni-4.5Al-0.5Ti 1 170 170 862 125 25 3 0-4 0 HRC Alloy 400 Ni-31Cu-2.5Fe 550 80 240 35 40 11 0-1 50 HB Alloy K-500 Ni-30Cu-2Fe-1.5Mn-2.7Al-0.6Ti 1100 160 79 0 115 20 300 HB Nickel-copper alloys Nickel-molybdenum and nickel-silicon alloys Hastelloy B Ni-28Mo-5.5Fe-2.5Co 834 Hastelloy D Ni-9.25Si-3Cu-1.5Co 386 56 63 92 HRB 85 345 50 10 93 HRB 79 3 Investment cast 121 586 Sheet 115 3 0-3 9 HRC Nickel-chromium-iron... Ti-5Al-2.5Sn 79 0 115 76 0 110 0.05 0.08 0.02 0.50 0 .20 5 2.5 Ti-8Al-1Mo-1V 900 130 830 120 0.05 0.08 0.015 0.30 0.12 8 1 1V Ti-2.25Al-11Sn5Zr-1Mo 1000 145 900 130 0.04 0.04 0.008 0.12 0. 17 2.25 11.0 5.0 1.0 0.2 Si Ti-6Al-2Sn-4Zr2Mo 900 130 830 120 0.05 0.05 0.0125 0.25 0.15 6.0 2.0 4.0 2.0 0.08 Si Ti-6Al-4V(A) 900 130 830 120 0.05 0.10 0.0125 0.30 0 .20 6.0 4.0V Ti-6Al-2Sn-4Zr6Mo(b) 1 170 170 ... Ni-15.5Cr-16Mo-5.5Fe-3 .75 W-1.25Co + V 78 5 114 372 54 62 209 HB Alloy 625 Ni-21.5Cr-9Mo-3.65Nb + Ta-2.5Fe 930 135 5 17 75 42.5 190 HB 690 100 320 47 50 79 HRB 690 100 310 45 45 Nickel-chromium-iron-molybdenum-copper alloys Hasteloy G Ni-22.25 Co,Nb,Ta Cr-19.5Fe-6.5Mo-2Cu Alloy 825 Ni-21.5Cr-30Fe-3Mo-2.25Cu + Al + The low-alloy nickels contain 94% min Ni The 5% Mn solid-solution addition in Nickel 211 protects against sulfur... Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 109 9- 1201 25 J.C Harkness, W.D Speigelberg, and W.R Cribb, Beryllium-Copper and Other Beryllium-Containing Alloys, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 40 3-4 27 26 W.L Mankins and S Lamb, Nickel and Nickel... Properties and Selection: Nonferrous Alloys and SpecialPurpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 82 2-8 39 29 D.W Dietrich, Magnetically Soft Materials, Properties and Selection: Nonferrous Alloys and SpecialPurpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 76 1 -7 81 30 E.L Frantz, Low-Expansion Alloys, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, ... Annealed 3 57 52 133 19 55 72 HRF Hard 532 77 441 64 8 82 HRB Annealed 378 55 119 17 45 80 HRF Half-hard 490 71 350 51 15 75 HRB Annealed 350 51 119 17 52 68 HRF Hard 420 61 318 46 7 80 HRB Annealed 350 51 175 25 55 40 HRB Pure copper C1 0200 OFHC 99.95 Cu C1 72 00 97. 9Cu-1.9Be-0.2Ni Co High-copper alloys Beryllium-copper or Brass C21000 Gilding, 95% Red brass, 85% Cartridge 70 % brass, C23000 C26000 C28000... and Nickel Alloys, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 42 8-4 45 27 J.J deBarbadillo and J.J Fischer, Dispersion-Strengthened Nickel-Base and Iron-Base Alloys, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 94 3-9 49 28 R.A Watson et al., Electrical... and strength The family of copper-nickel alloys also includes various dispersion- and precipitation-hardening alloys due to the formation of hardening phases with third elements, such as Ni2Si in C70250 (Cu-3Ni-0.7Si-0.15Mg) and the spinodal hardening obtainable in the Cu-Ni-Sn alloys (C7 270 0 with Cu-10Ni8Sn, for example) Copper-nickel-zinc alloys, also called nickel-silvers, are a family of solid-solution-strengthening... Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 109 9- 1201 19 D.E Tyler and W.T Black, Introduction to Copper and Copper Alloys, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 216240 20 D.E Tyler, Wrought Copper and Copper Alloy Products, Properties and Selection: . T 775 11 Extrusion 670 97 655 95 11 T6, T651 Sheet, plate 570 83 505 73 11 70 75 T73, T735X Plate, forging 505 73 435 63 13 T7351 Plate 505 73 435 63 15 74 75 T7651 Plate 455 66 390 57 15 (a). Ni 2 Si in C70250 (Cu-3Ni-0.7Si-0.15Mg) and the spinodal hardening obtainable in the Cu-Ni-Sn alloys (C7 270 0 with Cu-10Ni- 8Sn, for example). Copper-nickel-zinc alloys, also called nickel-silvers,. copper OFHC C1 0200 99.95 Cu . . . 22 1- 455 3 3- 66 6 9- 365 1 0- 53 5 5-4 . . . High-copper alloys Annealed 490 71 . . . . . . 35 60 HRB Beryllium-copper C1 72 00 97. 9Cu-1.9Be-0.2Ni or Co