Fig. 4 Schematic drawing of microstructure of tungsten sintered at different temperatures before being immersed into molten copper. (a) 1100 °C (2010 °F). (b) 1350 °C (2460 °F). (c) 1600 °C (2910 °F). (d) 2800 °C (5070 °F). 500×. Source: Ref 16 Infiltrated bodies that have a continuous skeleton phase above approximately 65 vol% can be shaped only by machining, while those having a smaller proportion of refractory metal dispersed as loose grains in the ductile metal matrix are plastically deformable at elevated temperatures (Ref 8). Soldering, brazing, or plating of the infiltrated product is aided by generally smooth surface films or high contents of infiltrant metal. The preceding binary systems illustrate combinations of two metals of widely differing melting temperatures that can be advantageously produced to near-net shape and full density by infiltration. Nickel also can be infiltrated into tungsten, but equilibrium at liquid-phase temperature causes severe attack of the refractory metal and requires careful process control to prevent incomplete penetration due to diffusion solidification. Coarse tungsten powder helps to produce compacts with larger capillary channels for better penetration. High heating rates, especially after about 90% of the absolute melting temperature of nickel is reached, improve infiltration conditions. If nickel is alloyed with copper, the solubility of tungsten in the liquid phase decreases as the copper percentage is increased. Infiltration is more practical (Ref 8) for heavy alloy systems because tungsten powder and process control requirements are less stringent. If nickel is alloyed with iron, however, the solid-liquid phase interaction is similar to the binary tungsten-nickel system, and infiltration again is more difficult. The same principle applies to the more complex refractory metal systems with nickel-chromium and cobalt-chromium alloys. Nevertheless, skeleton bodies of tungsten and molybdenum, as well as bodies of binary 85W-15Cr and 75W-25Cr alloys, can be successfully infiltrated with superalloys of the Nichrome-V, Hastelloy-C, Stellite, and Vitallium compositions into shapes simulating mechanical and engine test specimens (Ref 17). Several other refractory metal-based composite structures can be readily produced by infiltration. These include the high- density tungsten-lead system to produce materials suitable for shielding against radiation and the chromium-copper system to produce compositions for welding electrodes (Ref 8). To retain the low liquid-solid contact angles in these systems, a strong reducing atmosphere is necessary to prevent oxide films on the molten lead or on the solid chromium. A free metallic surface requires sintering of the skeleton above 1250 °C (2280 °F) to reduce any oxide film on the solid chromium. Carbide-Based Systems. Liquid-phase sintering of tungsten carbide/cobalt or titanium carbide/nickel systems capitalizes on the eutectics of the two phases. A limited solubility of the carbide in the matrix metal facilitates bonding; carbon and metal diffusion through the liquid are less important in densification than reactions at the carbide-metal phase boundaries. It is unknown whether the carbide particles in these systems form a rigid skeleton, but the interfacial tension between crystals of the carbide and the liquid metal appears to be anisotropic (Ref 18). During cooling from the sintering temperature, some or most of the carbon and metal dissolved in the liquid precipitates on grains that remained solid during the process. This mechanism can be altered somewhat if the rigid skeleton is formed first and the liquid phase subsequently infiltrates into the pore system. By first saturating the matrix metal with carbon and skeleton metal, the liquid phase dissolves less skeleton materials, and shape distortion is diminished. Carbide coalescence and grain growth also are decreased. The earliest attempts to produce cemented carbides were made by infiltrating carbide skeletons with unalloyed binder metals (Ref 19). Later, binder metal prealloyed with elements of the skeleton to inhibit contact face erosion was used in a broad investigation of infiltrating single and double carbides with many cobalt and nickel alloys (Ref 20). Table 2 lists some of the alloys used. The large number of feasible combinations includes several noteworthy successes, especially for titanium carbide-based stems with Nichrome and Vitallium infiltrants. A feasibility study (Ref 17) produced similar results for the same infiltration systems and laid the foundation for an extensive development program to utilize these materials for heat-resistant applications. Table 2 Carbide infiltration test matrix and evaluation Skeleton (a) Infiltration (b) Infiltrated product Temperature, Specimen No. Composition Density, g/cm 3 Pore volume, % Infiltrant composition, % Type °C °F Time, min Analysis, % Hardness, HRA Surface condition Microstructure Bench test (c) Specimen No. 1a 100Co 1500 2730 15 1a 1b 95Co-5WC 1460 2660 5 86-87 Contact face erosion, slight residue Very tough 1b 1c 75Co-25WC 1390 2530 5 1c 1d WC (6.1% C) 11.05- 13.35 29.3- 14.6 60Co-40WC Contact, opposing sides (d) 1350 2460 5 85-86 No erosion, heavy residue Porous in core, graphite precipitates, fairly uniform grain size Tough 1d 2a 100Co 1500 2730 15 2a 2b 80WC- 20TiC (e) 6.65- 8.34 37.1- 21.3 95Co-5WC Contact, opposing sides (d) 1460 2660 15 85.6-87 Contact face erosion Uniform phase distribution, some porosity Tough 2b 3a 80Co-20Cr 1500 2730 15 87.5-88 Slight erosion Very tough 3a 3b 66Co-28Cr- 6Mo 1450 2640 15 89.5-90 Slight porosity Fairly tough 3b 3c 72.7Co- 17.3Cr- 10TiC 1400 2550 15 88 No erosion, smooth Very tough 3c 3d TiC (18.8% C) 3.01- 3.63 33.2- 19.3 80Ni-20Cr Contact, opposing sides (d) 1450 2640 15 24.6Ni, 6.1Cr, bal TiC 83.5-85 Slight erosion Dense, uniform Very tough 3d 4a 80Co-20Cr 1500 2730 5 88+ Very tough 4a 4b 66Co-28Cr- 6Mo 1450 2640 5 89.5 Heavy contact face erosion and residue 4b 4c 72.7Co- 17.3Cr- 10TiC Contact, one side (d) 1400 2550 5 88+ Less contact face residue than in 4a Higher matrix concentration near contact face, porosity increasing toward far end Tough where dense, brittle where porous 4c 4d 97TiC- 3Mo 2 C (e) 3.38- 4.03 25.7- 11.4 80Ni-20Cr Capillary dip in molten infiltrant 1550 2820 3 22.5Ni, 5.7Cr, 2.1Mo 2 C, bal TiC 84.5-85 Alloy skin becoming heavier toward bottom, forming excess on bottom end Uniform phase distribution, generally dense Tough 4d 5a 80Co-20Cr 1500 2730 5 88+ 5a 5b 66Co-28Cr- 6Mo 1450 2640 5 90 Contact face erosion, some residue 5b 5c 95TiC- 5Mo 2 C (e) 3.46- 4.05 24.8- 11.6 72.7Co- Contact, one side (d) 1400 2550 5 88.5 Slight Similar to 4a-c Tough 5c 17.3Cr- 10TiC contact face residue 5d 72.7Ni- 17.3Cr- 10TiC Capillary dip in molten infiltrant 1550 2820 3 85 Similar to 4d Similar to 4d Tough 5d 6 90TiC- 10Mo 2 C (e) 3.54- 4.14 26.0- 14.6 80Ni-20Cr Contact, opposing sides (d) 1400 2550 15 22.9Ni, 5.5Cr, 7.1Mo 2 C, bal TiC 85-86 Slight contact face erosion, small residue, slightly porous Porous in core, less uniform phase distribution than in 4d Less tough than 4 and 5 6 7 70TiC- 30Mo 2 C (e) 4.09- 4.75 22.9-9.9 80Ni-20Cr Contact, opposing sides (d) 1400 2550 15 22.6Ni, 5.6Cr, 21.4Mo 2 C, bal TiC 86-87 Similar to 6, but more porous More porous in core, less uniform phase distribution than in 6 More brittle than 6 7 8 50TiC- 50Mo 2 C (e) 4.69- 5.58 21.3-8.3 80Ni-20Cr Contact, opposing sides (d) 1400 2550 15 22.3Ni, 5.7Cr, 35.8Mo 2 C bal TiC 86-87 Similar to 7, but more porous Very porous, nonuniform phase distribution More brittle than 7 8 Source: Ref 19 (a) All skeletons were presintered at 950 °C (1740 °F) and high sintered at 1500 °C (2730 °F) for 2 h in a carbon tube resistor furnace under hydrogen, except No. 1 to 3, which were high sintered in vacua in a carbon susceptor induction furnace. (b) In vacuum induction furnace. (c) A qualitative assessment of resistance against fragmentation by hammer blows. (d) Infiltrant mass was 40 to 45% of mass of infiltrated product. (e) Solid solution. The structure of infiltrated carbides reflects infiltration mechanics on a macroscale. In zones penetrated by the infiltrant, fully dense regions and slight expansion of the skeleton due to carbide grain separation are observed. Substantial porosity and some shrinkage occur in areas inadequately penetrated by the liquid alloy, such as the side opposite the contact face in unidirectional infiltration, or in the center for infiltration from opposite sides. Subsequent heat treatment is ineffective in eliminating porosity. Erosion at the contact faces is greatly diminished by infiltrating skeletons composed of tungsten- base multicarbides, so that saturating the infiltrant with skeleton elements by prealloys frequently is not required. The microstructure reflects the crystallographic characteristics of the carbide, and fully infiltrated regions do not differ in grain size and morphology from material whose liquid phase was sintered in situ. Rectangular and triangular grains are retained in infiltrated tungsten carbide, whereas for infiltrated titanium carbide, the cubic lattice is reflected by distinctly rounded grains. Grains of solid-solution carbides of tungsten and titanium or titanium and molybdenum are slightly rounded at the corners. Graphite precipitates accompany porosity in poorly infiltrated regions for tungsten carbide and titanium carbide skeletons, especially if starting powders contain more than a trace of free carbon. Ferrous-Base Systems. The thermodynamic affinity between solid iron and liquid copper offers the potential for virtually pore-free P/M products by infiltration. Moreover, the excellent wetting characteristics that exist in the brazing of steel can also be utilized for joining disparate bodies during infiltration. A powder compact and a casting or forging, or halves of complex or offset configurations, can be joined. Figure 5 (Ref 10) illustrates this self-brazing capacity without strength degradation. Finally, the generally smooth cupric film surrounding the infiltrated body serves as a base for surface coating or plating. Fig. 5 Infiltration-brazed butted iron-copper bars. Green bars were clamped together end-to- end, and one free end surface was contacted with molten infiltrant. Rupture of composite bar occurred away from the joint, evidence of the high strength of the brazed bond. (a) Bu tt joined bars after infiltration. (b) Machined tensile bar. (c) Tested tensile bar The ability of the infiltration process to combine major proportions of normally unalloyable industrial premier metals (iron and copper) was recognized as early as World War I (Ref 21, 22), but it was only in the late 1940s, through refinements in technique, that sound products could be made (Ref 10, 23, 24, 25). These products, in turn, have culminated in the present advanced state of the art. While copper content must be higher than for most commercially sintered iron-copper alloys because of the need to maintain an interconnected pore system for complete infiltration, good mechanical properties can be realized. This is apparent from the tensile strength-elongation data given in Fig. 6 for commercial iron powder with and without graphite additions (Ref 24). Strength is enhanced because the infiltrated structure, with a minimal amount of isolated pores, is virtually free of internal notches. The iron-copper system permits a precipitation-strengthening mechanism. If carbon is diffused into the iron to produce a hypoeutectoid structure of the skeleton, hardening by martensite transformation is possible. Where other metals are alloyed with the copper, solid solution strengthening of the matrix can be achieved. Fig. 6 Mechanical property ranges of copper- infiltrated iron and hypoeutectoid steel compacts before and after heat treatment In the binary iron-copper system, about 3.8 to 4% Fe is dissolved by the liquid copper under equilibrium conditions at an infiltration temperature of 25 to 50 °C (45 to 90 °F) above the peritectic temperature of 1090 °C (2000 °F) at the copper side of the phase diagram, while the -iron dissolves about 8 to 8.5% of the copper. At 900 °C (1650 °F), the solubility of -iron in copper is about 1.5%; it diminishes to less than 0.04% at room temperature, at which point the solubility of copper in iron is equally low. These thermodynamic relations form the basis for precipitation hardening. However, macrodispersion of the two phases in the infiltrated alloys causes concentration of precipitates in thin zones at the phase boundaries, as shown in Fig. 7 (Ref 26). Consequently, conventional hardness tests are not precise enough to show a noticeable increase in macrohardness after a precipitation treatment, such as quenching from 900 °C (1650 °F) followed by prolonged tempering at 600 °C (1110 °F). The precipitation mechanism produces increases in strength, elongation, and impact resistance (Ref 8) and also can be traced through changes in the electrical conductivity (Ref 25). Fig. 7 Microstructure of 25Cu-Fe compact sintered at 1100 °C (2010 °F) for 30 min. Large, rounded, dark areas are -iron, separated by a diffusion layer from the light copper phase containing the fine, dark, iron- rich precipitate. 1000×. Source: Ref 26 If carbon is diffused into the iron skeleton with copper infiltrated afterward, two processes compete with one another during cooling. During quenching, normal martensitic transformation occurs inside the skeleton structure. At the same time, however, precipitation of the dissolved iron in the copper and copper in iron, respectively, is suppressed in the phase boundary zones. During reheating, martensite decomposition causes a decrease in hardness of the steel skeleton structure, with simultaneous increases in hardness and strength of the boundary zones. Table 3 shows the effect of such a heat treatment on the mechanical properties of a copper-infiltrated 0.3% carbon steel. Water quenching produces an appreciable increase in hardness and brittleness. With increasing reheating temperatures, however, the material becomes softer and tougher, without loss in strength (Ref 8). Table 3 Effect of heat treatment on the mechanical properties of 0.3% carbon steel infiltrated with 11 vol% Cu Ultimate tensile strength Impact resistance (b) Treatment (a) Hardness, HV MPa ksi Elongation, % J ft · lb Copper-infiltrated and furnace cooled 262 704 102.1 7.3 22.16 16.35 Reheated to 900 °C (1650 °F) and water quenched 437 790 114.6 5.2 11.76 8.68 Reheated for 2 h at 400 °C (750 °F) 360 . . . . . . . . . 14.71 10.85 Reheated for 2 h at 500 °C (930 °F) 302 771 111.8 7.6 23.14 17.07 Reheated for 2 h at 600 °C (1110 °F) 255 750 108.8 13.5 62.75 46.30 (a) Hametag iron powder pressed to 88% of theoretical density, sintered at 1220 °C (2230 °F) for 1 h in hydrogen, and infiltrated with electrolytic copper at 1100 °C (2010 °F) for 30 min. (b) Unnotched test bar of 1 cm 2 cross section. When a copper alloy is used as an infiltrant, other benefits can accrue. When precipitates form in the matrix, such as copper alloys with beryllium, chromium, or silicon, the resulting strength increase during heat treating of the infiltrated body reinforces iron-copper precipitation zones at the boundaries. The strengthening effect of precipitates is independent of heat treatment and augments strength increases obtained with unalloyed copper infiltrant in solid-solution alloys. However, some solid-solution alloy infiltrants minimize erosion at the point of initial contact with the iron skeleton, because these alloys melt slowly over a range of temperatures. Brass containing 20% Zn exhibits this phenomenon. A copper alloy with manganese in amounts up to 5%, especially if it also contains sufficient iron to inhibit a severe attack of the skeleton contact faces, has a different beneficial effect. The manganese oxidizes preferentially in the commercial atmospheres where infiltration takes place. A nonadhering porous crust results that can be removed during finishing more easily than the tenacious residue formed by a binary copper-iron infiltrant (Ref 27, 28). Many ferrous metal-base infiltration systems have been explored experimentally. Kieffer and Benesovsky (Ref 8) have investigated the iron-gold, iron-bismuth, iron-cadmium, iron-lead, iron-antimony, and iron-tin systems for bearings, and the iron-cobalt-silicon, iron-copper-silicon, and iron-manganese-silicon systems for magnetic or structural parts. Alloys of the iron-zinc system have also been produced by infiltration, but treatment in a pressure vessel is required to overcome the high vapor pressure of the zinc (Ref 29). Austenitic stainless steel skeletons infiltrated with silver possess excellent corrosion resistance, thus making them suitable for food processing applications. Ferritic stainless steel and high-manganese steel compacts with varying carbon contents also display improved corrosion resistance when infiltrated with cupric alloys. These alloys also exhibit extraordinary hardness and wear resistance, coupled with a considerable toughness. Table 4 lists mechanical and technological properties of several ferrous-base infiltrated materials. Table 4 Properties of some ferrous metal-based infiltrated alloys Density, g/cm 3 Ultimate tensile strength Impact resistance (a) Skeleton composition, % Infiltrant composition, % Infiltrant, vol% Calculated Determined Hardness, HV MPa ksi Elongation, % J ft · lb Comment 100Fe 100Pb 10 8.15 7.95 93.5 251 36.4 14 . . . . . . Free machining, extrudable 99Fe-Cu 100Ag 11 8.09 8.00 178 378.5 54.9 11 . . . . . . . . . 100Fe 80Cu-20Ni 17 8.04 7.76 213 419.5 60.9 8 21.57 15.91 . . . 100Fe 65Cu-35Mn 13 7.92 7.63 256 446 64.7 10 31.37 23.14 . . . 93.2Fe-6Mn-0.8C 100Cu 13 7.90 7.87 740 . . . . . . . . . . . . . . . Naturally hard, wear resistant 87.2Fe-12Mn-0.8C 100Cu 9 7.89 7.69 310 562 81.5 6 . . . . . . Wear resistant, work hardening 93.5Fe-3Cr-3Mn-0.5C 100Cu 14 7.96 7.93 502 957 138.8 4 . . . . . . . . . (a) Unnotched test bar of 1 cm 2 cross section. Nonferrous-Based Systems. The major nonferrous metals with higher melting points are thermodynamically compatible with many low-melting metals in the liquid state. Consequently, skeleton bodies of cobalt and nickel can be readily infiltrated with gold, as well as with many of the low-melting heavy metals, such as bismuth, lead, or antimony (Ref 8). Copper can also be infiltrated into cobalt and nickel skeletons; however, because of the formation of solid solution in all proportions, infiltration into nickel powder compacts requires a narrow particle size range, wide capillaries within a pore volume not exceeding about 35%, short infiltration time, and a vacuum to assist the capillary forces. Mercury wets nickel well without forming an amalgam and is easily impregnated into skeleton bodies, provided the pore structure prevents exudation of the heavy liquid metal (Ref 30). Copper is another skeleton metal with pores that can be readily filled with liquid low-melting-point metals, such as lead (Ref 31, 32) or bismuth (Ref 33). Vacuum impregnation is suitable for incorporating lead-base alloys, such as those containing 15% Sb and 5 to 10% Sn, into spongy structures of nickel-copper or nickel-iron supported by steel backing (Ref 34). Combinations of aluminum or aluminum alloys with low-melting metals such as bismuth, lead, thallium, or thallium-lead by means of infiltration in hydrogen or a vacuum have been proposed (Ref 33, 35). However, strict control of powder characteristics, especially particle shape and surface condition, to maximize wetting appears to put infiltration at a disadvantage over in situ liquid-phase sintering. To overcome this problem, zinc or cadmium can be added to the aluminum of the skeleton, followed by cleaning and activating the free surface of the pores by evaporation of the lower boiling metal before proceeding with the impregnation of a metal such as lead (Ref 36). References cited in this section 8. R. Kieffer and F. Benesovsky, The Production and Properties of Novel Sintered Alloys (Infiltrated Alloys), Berg- und Hüttenmännische Monatshefte, Vol 94 (No. 8/9), 1949, p 284-294 10. F.V. Lend, Powder Metallurgy, Metal Powder Industries Federation, 1980, p 313-319 15. C.G. Goetzel, Treatise on Powder Metallurgy, Vol 2, Interscience, 1950, p 196 16. K. Schröter, Border Regions of Metallography, Z. Metallkd., Vol 23 (No. 7), 1931, p 197-201 17. J.M. Kro l and C.G. Goetzel, "Refractory Metal Reinforced Super Alloys," USAF Technical Report 5892, ATI No. 57154, May 1949 18. F.V. Lend, Powder Metallurgy, Metal Powder Industries Federation, 1980, p 383-400 19. R. Kieffer and F. Kölbl, Production of Hard Metals by Infiltration, Berg. Hüttenmämann. Monatsh., Vol 95 (No. 3), 1950, p 49-58 20. H. Baumhauer, Hard Tools and Process for Making Them, U.S. Patent 1,512,191, 1924, German Patent 443,911, 1927 21. L. Reimann, Production of Bodies from Metallic Compound s, German Patent 300,669, 1917; L. Reimann and H. Leiser, Metallic Alloy, British Patent 148,533, 1921 22. C.L. Gebauer, Process of Producing Metal Bodies, U.S. Patent 1,342,801, 1920; C.L. Gebauer, Production of a Composite Metallic Article, U.S. Patent 1,395,269, 1921 23. F.P. Peters, Cemented Steels A New High-Strength Powder Metallurgy Product, Materials and Methods, Vol 23 (No. 4), 1946, p 987-991 24. E.S. Kopecki, Cemented Steels, Iron Age, Vol 157 (No. 18), 1946, p 50-54 25. C.G. Goetzel, Cemented Steels Infiltration Studies with Pure Iron and Copper Powders, Powder Metall. Bull., Vol 1 (No. 3), 1946, p 37-43 26. L. Northcott and C.J. Leadbeater, Sintered Iron-Copper Compacts, Symposium on Powder Metallurgy, Special Report No. 38, The Iron and Steel Institute, 1947, p 142-150 27. P. Schwarzkopf, Infiltration of Powder Metal Compacts with Liquid Metal, Met. Prog., Vol 57 (No. 1), 1950, p 64-68 28. G. Stern, The Effect of Infiltration on Physical Properties of Sinterings, Prec. Met. Mold., Vol 1 1 (No. 6), 1953, p 92-102 29. J. Schramm and A. Mohrnheim, Precipitation Hardening of Iron-Zinc and Cobalt-Zinc Alloys, Z. Metallkd., [...]... 10. 29 10.30 9. 30 9. 60 Hardness, HB 23 0-2 50 21 0-2 30 20 0-2 20 16 0-1 80 14 0-1 60 19 0-2 10 18 0-1 90 16 0-1 70 14 0-1 50 12 0-1 30 (b) (c) Transverse rupture strength MPa ksi 124 0-1 400 18 0-2 00 110 0-1 240 16 0-1 80 96 5-1 100 14 0-1 60 86 0 -9 65 12 5-1 40 79 0-8 25 11 5-1 20 110 0-1 240 16 0-1 80 96 5-1 100 14 0-1 60 86 0 -9 65 12 5-1 40 52 5-8 60 11 0-1 25 69 0-5 25 10 0-1 10 450 65 350 50 Electrical conductivity Mmho/cm %IACS 0.2 0-0 .23 3 5-4 0 0.2 4-0 .25... 42, 199 4, p 138 9- 1 397 78 T.R Jonas, J.A Cornie, and K.C Russell, Infiltration and Wetting of Alumina Particulate by Aluminum and Aluminum-Magnesium Alloys, Metall Trans A, Vol 26, 199 5, p 1 49 1-1 497 79 S Long, Z Zhang, and H.M Flower, Characterization of Liquid Metal Infiltration of a Chopped Fiber Preform Aided by External Pressure, Parts I-III, Acta Mater., Vol 43, 199 5, p 348 9- 3 5 09; Vol 44, 199 6,... and W Hotop, Powder Metallurgy and Sintered Materials, Springer-Verlag, 194 3, p 32 4-3 26, 3 29 38 F.R Hensel, E.I Larsen, and E.F Swazy, Physical Properties of Metal Compositions with a Refractory Metal Base, Powder Metall., J Wulff, Ed., American Society for Metals, 194 2, p 48 3-4 92 39 C.G Goetzel, Treatise on Powder Metallurgy, Vol 2, Interscience, 195 0, p 20 7- 2 09, 21 6-2 17 40 P Schwarzkopf, Powder Metallurgy,... Pressure, Parts I-III, Acta Mater., Vol 43, 199 5, p 348 9- 3 5 09; Vol 44, 199 6, p 423 3-4 240 80 T Yamauchi and Y Nishida, Consideration on Suitable Infiltration Conditions for Molten Metal into Fibrous Preforms, J Jpn Inst Light Met., Vol 45, 199 5, p 40 9- 4 14 81 J.-H Ahn, N Terao, and A Berghezan, Experimental Factors on Wetting and Infiltration Fronts in Metals, Scr Mater., Vol 22, 198 8, p 79 3-7 96 82 J.-H Ahn and. .. Springer-Verlag, 194 3, p 32 4-3 26, 3 29 38 F.R Hensel, E.I Larsen, and E.F Swazy, Physical Properties of Metal Compositions with a Refractory Metal Base, Powder Metall., J Wulff, Ed., American Society for Metals, 194 2, p 48 3-4 92 39 C.G Goetzel, Treatise on Powder Metallurgy, Vol 2, Interscience, 195 0, p 20 7- 2 09, 21 6-2 17 40 P Schwarzkopf, Powder Metallurgy, Macmillan, 194 7, p 16 7-1 68 41 C.G Goetzel, Treatise on Powder. .. in Powder Metallurgy, Proc of 15th Annual Meeting, Metal Powder Industries Federation, 195 9, p 5 6-6 6 8 R Kieffer and F Benesovsky, The Production and Properties of Novel Sintered Alloys (Infiltrated Alloys), Berg- und Hüttenmännische Monatshefte, Vol 94 (No 8 /9) , 194 9, p 28 4-2 94 34 A.L Boegehold, Copper-Nickel-Lead Bearings, Powder Metall., J Wulff, Ed., American Society for Metals, 194 2, p 52 0-5 29. .. Alloy, Part I: Theory, Metall Trans A, Vol 21, 199 0, p 205 9- 2072 58 V.J Michaud and A Mortensen, Infiltration of Fiber Preforms by a Binary Alloy, Part II: Further Theory and Experiments, Metall Trans A, Vol 23, 199 2, p 226 3-2 280 59 R.B Calhoun and A Mortensen, Infiltration of Fibrous Preforms by a Pure Metal, Part IV: Morphological Stability of the Remelting Front, Metall Trans A, Vol 23, 199 2, p 2 29 1-2 299 ... Metal, Part III: Capillary Phenomena, Metall Trans A, Vol 21, 199 0, p 225 7-2 263 57 A Mortensen and V.J Michaud, Infiltration of Fiber Preforms by a Binary Alloy, Part I: Theory, Metall Trans A, Vol 21, 199 0, p 205 9- 2072 58 V.J Michaud and A Mortensen, Infiltration of Fiber Preforms by a Binary Alloy, Part II: Further Theory and Experiments, Metall Trans A, Vol 23, 199 2, p 226 3-2 280 59 R.B Calhoun and. .. on Powder Metallurgy, Vol 2, Interscience, 195 0, p 196 16 K Schröter, Border Regions of Metallography, Z Metallkd., Vol 23 (No 7), 193 1, p 19 7-2 01 17 J.M Krol and C.G Goetzel, "Refractory Metal Reinforced Super Alloys," USAF Technical Report 5 892 , ATI No 57154, May 194 9 18 F.V Lend, Powder Metallurgy, Metal Powder Industries Federation, 198 0, p 38 3-4 00 19 R Kieffer and F Kölbl, Production of Hard Metals... Yield strength MPa ksi FX-1005 AS 0. 3-0 .6 570 83 440 HT 0. 3-0 .6 830 120 AS 0. 6-1 .0 620 HT 0. 6-1 .0 FX-2000(d) AS 0.3 max FX-2005(e) AS 0. 3-0 .6 HG 0. 3-0 .6 AS 0. 6-1 .0 HG 0. 6-1 .0 FX-1008 FX-2008(f) 8.014 .9 8.014 .9 8.014 .9 8.014 .9 15.025.0 15.025.0 15.025.0 15.025.0 15.025.0 80. 591 .7 80. 591 .7 80. 191 .4 80. 191 .4 70.785.0 70.484.7 70.484.7 70.084.4 70.084.4 Impact energy(c) J ft · lb 19 14 Apparent hardness . 19 0-2 10 110 0-1 240 16 0-1 80 0.2 8-0 .30 4 9- 5 3 500 25 . . . 75 . . . 10.27 18 0-1 90 96 5-1 100 14 0-1 60 0.3 0-0 .32 5 3-5 6 400 30 . . . 70 . . . 10.28 16 0-1 70 86 0 -9 65 12 5-1 40 0.3 2-0 .34 5 6-6 0. 0. 3-0 .6 8. 0- 14 .9 80. 5- 91 .7 570 83 440 64 4.0 19 14 75 HRB 135 20 FX-1005 HT 0. 3-0 .6 8. 0- 14 .9 80. 5- 91 .7 830 120 740 107 1.0 9. 5 7.0 35 HRC 135 20 AS 0. 6-1 .0 8. 0- 14 .9 80. 1- 91 .4. 20 0-2 20 96 5-1 100 14 0-1 60 0.2 6-0 .28 4 5-4 9 500 25 75 . . . . . . 15.8 16 0-1 80 86 0 -9 65 12 5-1 40 0.2 8-0 .30 4 9- 5 3 350 30 70 . . . . . . 15.2 14 0-1 60 79 0-8 25 11 5-1 20 0.3 0-0 .33 5 3-5 7 250 20 .