Volume 07 - Powder Metal Technologies and Applications Part 16 doc

Fatigue and Fracture Control for Powder Metallurgy Components * Randall M. German and Richard A. Queeney, The Pennsylvania State University Porosity Effects An inherent physical characteristic of P/M materials is the presence of pores. The role of porosity in determining fatigue endurance in powdered metals is akin to that of porosity that is induced through metal solidification in casting or welding. However, the porosity that is characteristic of sintered powdered metals, and of those materials subsequently deformation processed, may differ in character and influence from solidification porosity. In all cases, porosity catches the attention of the design engineer because it immediately conjures up images of classical stress concentrators. In addition, porosity featuring sharp re-entrant corners, a possibility for marginally equilibrium-sintered powder particle boundaries, may be more accurately viewed as crack precursors. Pore geometry can be altered by modifications to the sintering cycle, such as a longer hold time or higher temperature, wherein the smoother pores improve strength, fatigue life, and fracture resistance. Ductility is sensitive to the pore structure, but impact, fracture, and fatigue behavior have the greatest sensitivities. In general, dynamic strength responses are the most sensitive. Even in those materials possessing full density, inferior properties can occur due to microstructural defects. Recent applications have pushed P/M into very demanding applications, a good example being the automotive connecting rod. Such rods are formed by hot forging a porous P/M preform using an alloy of Fe-2Cu- 0.8C. They weigh nearly 650 g. The ultimate tensile strength is 825 MPa (120 ksi), with a yield strength of 550 MPa (80 ksi) and fatigue endurance limit of 255 MPa (37 ksi). With expansion of P/M fabrication technology into dynamically loaded components, there arise performance limitations associated with fracture and fatigue. For static tensile properties, there is a good basis for predicting the effect of residual porosity on strength (Ref 2, 3, 4, 5). Recent research has had an emphasis on fatigue and fracture behavior to fill the database void. Porosity is especially important to the high-cycle fatigue life (Ref 6, 7, 8, 9). Pores play a role in both crack initiation and propagation, typically increasing the threshold stress intensity for crack initiation but lowering the resistance to crack propagation (Ref 10). For an alloy of Fe-2Ni-0.8C the fatigue endurance limit at 10 7 reverse cycles is between 200 and 250 MPa at a density of 7.1 g/cm 3 (approximately 10% porosity). That is approximately 35% of the tensile strength, and many porous injection-molded materials exhibit similar ratios of fatigue strength to tensile strength. Closed pores are less detrimental than those that are interconnected and open to the external surface. Surface pores act as the preferred site of fatigue crack initiation by acting as stress concentrators (Ref 11). The simplest view of a pore would be that of a spherical hole embedded in a continuous matrix whose response parameters are identical with that of a fully dense form of exactly the same metallurgical state. Since the maximum, most positive, principal stress is the controlling load parameter for high-cycle fatigue endurance, the embedded pore in a tensile field is a reasonably approximate model for a stressed porous metal without interacting stress fields (one of less than 10% porosity). The local pore-dominated stress field can be expected to control local events such as crack initiation and threshold stress intensity values ( K th ). For an embedded pore, the local stress fields are related to the far-field (applied) design stresses (Ref 12): (Eq 1) Note that, for the average steel, v = 0.30, and the local stresses are magnified by a factor of about 2. The concentrated stress field does not persist far beyond the embedded pore, being reduced to 105% of the design value at a distance of 2a (where a is the pore radius) from the pore center, thus possibly exerting little influence on far field parameters. If the pore features sharp re-entrant corners, the elastic concentrated stresses are more accurately predicted from an elliptical hole example. However, these same pores are unlikely to exhibit smooth ellipsoidal morphologies, and their concentration effects can be more usefully predicted by the proportionality (Ref 13): (Eq 2) Here, a is the major pore length normal to the maximum principal stress direction, and is the radius of curvature at the sharp corner. Suffice to say that the calculated stress concentration value can be large for a nonspheroidized pore. Due to the stress-raising properties of pores, and the fact that most fatigue failures originate at free surfaces, treatments aimed at surface densification and serendipitous surface strengthening (e.g., coining, shot peening, ausrolling) raise the fatigue endurance limit (Ref 14, 15, 16). As a consequence, the current models for fatigue response in porous sintered materials have a major dependence on the pore microstructure: the models address the total porosity, alloying homogeneity in near-pore regions, pore size, pore shape, and interpore separation distance (Ref 6, 17, 18, 19, 20, 21, 22, 23, 24). Round pores provide improved resistance to crack propagation. Pores act as linkage sites through which cracks can propagate. Microstructure- based fatigue models for ferrous alloys have had to address seemingly contradictory porosity effects: round pores retard stable fatigue crack propagation but increase crack extension growth rates by contributing linkage sites. While extant theories successfully explain some porosity effects on crack propagation, no total predictive model has been created that embraces microstructure effects (Ref 18, 20, 22). The presence of pores in the reversed plastic zone that is the site of propagating crack damage does not lend itself to facile analysis. Pore structure changes are obtainable through processing and material variations: powder variables (particle size distribution, particle shape); compaction variables (type of lubricant, amount of lubricant, tool motions, maximum pressure); sintering variables (hold time, maximum temperature, atmosphere); and postsintering treatments. Figure 1 compares the pore structure in two sintered stainless steels to emphasize this point. Smaller particles result in faster sintering and higher strengths and toughness. Associated with the smaller particle sizes are smaller final pores. At lower densities (around 6.6 g/cm 3 or 18% porosity in steels), a high-relative content of small particles is beneficial to fatigue resistance, while at higher densities (over 7.1 g/cm 3 or less than 10% porosity), larger particles prove beneficial. The difference relates to the ligament size between pores, which is the determinant of fatigue. There are three pore microstructure parameters relevant to the fatigue resistance of porous P/M materials: pore size, pore curvature, and pore spacing. These largely reflect the role of stress concentration with respect to the advancing fatigue crack, as noted above (Ref 6). Thus, lower porosity contents, smoother (rounder) pores, and wider interpore separations increase the fatigue endurance strength. Fig. 1 Two stainless steels fabricated by P/M, as demonstrations of the microstructure variations possible by tailoring the powder, compaction, and sintering variables. (a) A high- porosity microstructure useful for filtration, formed by press and sinter. 1000×. (b) A closed-porosity, high- density microstructure useful for mechanical components, formed by injection molding and high-temperature sintering. 200× Fatigue cracks have been successfully analyzed with regard to their propagation response, and the cyclic growth of a crack can be predicted by the modified Paris growth law: (Eq 3) where the material parameters A and n must be experimentally determined for any given material microstate, including different distributions of porosity. The stress analytical variable K eff , the effective stress intensity range, factors out that portion of the total stress range that relieves the stresses holding the crack flanks closed, the opening stress range op , leaving only the stress range component that displaces the crack faces relative to each other. In the presence of appreciable levels of mean stress, the relation between load design parameters and fatigue crack propagation rates is given by: (Eq 4) Here, the total stress-intensity factor range K is the load variable, but the fatigue ratio R = min / max and the fracture toughness K c , or K Ic for a low-ductility sintered member, enter into the fatigue crack propagation response, as do the material constants C and n. Regardless of which of the three fatigue crack propagation relations are relevant in a particular service context, their collective utility lies in predicting member lifetimes, or precise segments of that lifetime. Thus, in the case of the Paris law: (Eq 5) The calculated endurance cycles to failure, N, can be from an initial flaw size a 0 that may be the minimum detectable to final fracture at a = a c , where: K Ic = max (Eq 6) Again, the fracture toughness of the material plays a role in determining structural endurance. The fracture resistance K Ic , or K R in the case of more ductile sintered materials (Ref 25), is known to be porosity sensitive (Ref 26). As a first estimate, the sensitivity is about a 100 MPa gain in toughness per percentage point of porosity reduction in quenched and tempered steels. Copper-infiltrated steels have toughnesses that run from 40 to 50% those of wrought steels (Ref 27, 28). However, their static strengths are equivalent, reflecting the inability of the steel skeleton to absorb the same level of strain energy release as a fully dense body of the same material. The material parameters in the modified Paris equation (Eq 3) are sensitive to the porosity state, a not unexpected result since the coefficient and exponent are related to the plastic zone size r p ahead of the advancing crack, the only region of irreversible deformation. The coefficient A in Eq 3 increases with increasing porosity fraction (Ref 29), resulting in faster growth rates for comparable K eff values. With higher coefficient measures but constant exponent n values, the threshold stress intensity range K th , below which no crack growth is thought to occur, also decreases with porosity increase. However, K th values for fully dense materials are sufficiently low that it is not at all clear in what way they could be successfully employed in design practice if service stresses are to be set at appreciable fractions of the yield or tensile strength. The reversed plastic zone size is given by (Ref 30): (Eq 7) When the plastic zone size is of the same size as the average pore diameter, it is effectively enlarged by the high strain field in the vicinity of the pore. The net result of the enlarged effective plastic zone is reflected in higher values of the exponent n (Ref 11), and the enlarged zone is even more pronounced in the response of short cracks driven by the locally raised stress/strain fields associated with design stress concentrators (Ref 15). Copper-infiltrated steels (Ref 27, 28) are as fatigue and fracture resistant as fully dense wrought medium-strength steels. Although the performance standards are not up to those of fully dense martensitic steels, copper infiltration represents a considerable cost savings over forging to full density, while maintaining the cost advantage inherent in press and sinter P/M. Typical fatigue endurance limits (or fatigue strengths at about 10 7 cycles) are collected in Table 8 for several P/M materials. This compilation includes several ferrous alloys, reflecting the high interest in P/M fatigue for automotive applications. There are variations in density, alloying, and sintering cycles to show the relative effects on sintered properties. Systematic testing of various alloys has shown that slight changes in the particle size distribution or alloying homogeneity can affect these properties. For example, in the Fe-2Ni-0.8C alloy system, shifts in just the iron powder source lead to ±33 MPa (±4.8 ksi) variations in the fatigue endurance strength. Accordingly, the values in Table 8 are for relative ranking purposes only and cannot be used as an accurate basis for design of fatigue-sensitive components. Table 8 Representative P/M materials, processing cycles, and fatigue endurance limit Composition, wt% Processing Density, g/cm 3 Testing Endurance limit, MPa Al-5Cu-0.5Mg-0.8Si P+S, 600 °C, 1 h 2.6 . . . 53 Fe P+S, 1120 °C, h 6.0 Rotating R = -1 39 Fe P+S, 1120 °C, h 6.7 Bending R = 0 67 Fe P+S, 1120 °C, h 6.9 Rotating R = -1 102 Fe P+S, 1150 °C, 1 h 7.2 Axial R = -1 65 Fe P+S, 1120 °C, h 7.3 Rotating R = -1 145 Fe P+S, 1250 °C, 2 h 7.6 Rotating R = -1 181 Fe HIP, 1100 °C, 200 MPa 7.86 Axial R = -1 230 Fe-17Cr-4Cu-4Ni (17-4 PH) PIM, 1350 °C, 2 h, HT 7.5 Rotating R = -1 517 Fe-1.5Cu-0.6C P+S, 1120 °C, h, HT 7.0 Bending R = 0 390 Fe-2Cu-0.5C P+S, 1120 °C, 30 min 7.1 Rotating R = -1 125 Fe-2Cu-0.8C P+S, 1120 °C, h 6.7 Rotating R = -1 165 Fe-2Cu-0.8C P+S, 1120 °C, h 7.0 Rotating R = -1 234 Fe-2Cu-0.8C P+S, 1120 °C, h 7.15 Rotating R = -1 241 Fe-2Cu-0.8C P+S, 1330 °C, 1 h 7.1 Rotating R = -1 270 Fe-2Cu-2Ni-0.8C P+S, 1120 °C, h, HT 7.0 Rotating R = -1 240 Fe-2Cu-2Ni-0.8C P+S, 1120 °C, h, HT, SP 7.0 Rotating R = -1 282 Fe-2Ni-0.5C PIM, 1250 °C, 4 h, HT 7.7 Rotating R = -1 239 Fe-2Ni-0.8C P+S, 1120 °C, h, HT 7.12 Rotating R = -1 159 Fe-2Ni-0.8C P+S, 1175 °C, h 6.9 Rotating R = -1 192 Fe-2Ni-0.5Mo-0.5C P+S, 1120 °C, h, HT 7.0 Rotating R = -1 350 Fe-2Ni-0.5Mo-0.5C DP + DS, 1260 °C, 30 min, HT 7.4 Rotating R = -1 425 Fe-2Ni-0.5Mo-0.5C P+S, 1120 °C, 30 min, HT 6.8 Rotating R = -1 345 Fe-2Ni-0.5Mo-0.4C PF, 1150 °C 7.9 Rotating R = -1 780 Fe-2Ni-1Mo-0.9C P+S, 1275 °C, 1 h, HT . . . Rotating R = -1 390 Fe-2Ni-1.5Cu-0.5Mo-0.5C HIP, 1160 °C, 3 h, 105 MPa, HT 7.86 Rotating R = -1 480 Rotating R = -1 129 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1120 °C, h 7.1 Bending R = 0 148 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1120 °C, 2 h 7.1 Rotating R = -1 135 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1250 °C, h 7.1 Rotating R = -1 147 Rotating R = -1 195 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1250 °C, 2 h 7.1 Bending R = 0 266 Fe-7Ni PIM, 1250 °C, 1 h 7.71 Rotating R = -1 236 Tool steel (Fe-8Co-6.3W-5Mo-3V-4Cr-1.3C) HIP, 1150 °C, hot roll, HT 8.0 Axial R = -1 950 Ni P+S, 1000 °C, 1 h 7.8 Rotating R = -1 70 Ni DP + DS, 1300 °C, 3 h 8.5 Rotating R = -1 121 Ni 3 Si Hot extrude, HT . . . Axial R = -1 579 Ti-6Al-4V HIP, 925 °C, 3 h, 200 MPa 4.46 Axial R = -1 475 P+S, pressed and sintered; PIM, powder injection molded and sintered; HIP, hot isostatically pressed; HT, heat treated; SP, shot peened; DP + DS, double pressed and double sintered; PF, powder forged Fracture studies of P/M alloys are usually restricted to impact testing, and often this is performed in the unnotched condition because of the low toughness of porous materials. Fracture toughness measurements on P/M materials are relatively rare. Tables 9 and 10 summarize some prior findings. Table 9 demonstrates the density effect on tensile properties and K Ic , for an Fe-Ni-Mo-C steel. Note that as the porosity decreases, the strength essentially doubles, ductility increases substantially (a nearly tenfold gain), the impact energy goes up by a factor of 4, and fracture toughness increases threefold. Table 10 collects several examples of the mechanical properties of ferrous alloys and one titanium alloy as representative values obtainable via P/M. Clearly, porosity negatively affects the fracture toughness in most materials. The fracture toughness of P/M steels is essentially a linear function of density (Ref 18), with sensitivities of about 100 MPa gain in toughness per percentage point of porosity reduction. In low-density P/M materials, the fracture crack propagation is rapid because the pore structure amplifies the stress and provides an easy path. At low porosity levels, an advancing fracture crack can be blunted by the pores, effectively forming microcracks that improve toughness (Ref 31). Table 9 Mechanical properties of pressed and sintered Fe-1.8Ni-0.5Mo-0.5C P/M compacts Double pressed and double sintered, 1120 °C, h, tested as-sintered Porosity, % Density, g/cm 3 Elastic modulus, GPa Yield strength, MPa Ultimate strength, MPa Elongation, % Notched impact energy, J Fracture toughness MPa 16 6.6 110 280 350 2 3 19 10 7.1 145 370 460 3 4 28 5 7.4 180 425 610 5 4 38 0 7.9 190 590 800 19 12 65 Table 10 Representative P/M materials, processing cycles, and fracture toughness Composition, wt% Processing Density, g/cm 3 Fracture toughness, MPa Fe-4.4Cr-9.2Co-7.2V-3.7Mo-9.2W-2.7C P+S, 1150 °C, 1 h 8.1 13 Fe-1.5Cu-2Ni-0.8C P+S, 1120 °C, h 6.8 40 Fe-1.8Ni-0.5Mo-1.5C P+S, 1120 °C, h 6.6 15 Fe-1.8Ni-0.5Mo-1.5C P+S, 1150 °C, h 6.8 26 Fe-1.8Ni-0.5Mo-1.5C P+S, 1120 °C, h 7.1 24 Fe-1.8Ni-0.5Mo-1.5C DP + DS, 1100 °C, h 7.5 21-38 Fe-1.8Ni-0.5Mo-1.5C HF, 1100 °C 7.85 64 Fe-0.8P-0.3C P+S, 1120 °C, h 7.0 22 Fe-0.8P-0.3C DP + DS, 1120 °C, h 7.8 20 Ti-6Al-4V HIP, 925 °C, 3 h, 200 MPa 4.46 65 P+S, pressed and sintered; HIP, hot isostatically pressed; DP + DS, double pressed and double sintered; HF, hot forged References cited in this section 2. B. Karlsson and I. Bertilsson, Mechanical Properties of Sintered Steels, Scand. J. Metall., Vol 11, 1982, p 267- 275 3. G.F. Bocchini, The Influences of Porosity on the Characteristics of Sintered Materials, Rev. Powder Met. Phys. Ceram., Vol 2, 1985, p 313-359 4. R. Haynes, The Mechanical Behavior of Sintered Metals, Rev. Deform. Behav. Mater., Vol 3, 1981, p 1-101 5. S.H. Danninger, G. Jangg, B. Weiss, and R. Stickler, Microstructure and Mechanical Pr operties of Sintered Iron, Part 1: Basic Considerations and Review of Literature, Powder Met. Inter., Vol 25, 1993, p 111-117 6. B. Weiss, R. Stickler, and H. Sychra, High-Cycle Fatigue Behaviour of Iron-Base PM Materials, Metal Powder Report, Vol 45, 1990, p 187-192 7. R. Haynes, Fatigue Behaviour of Sintered Metals and Alloys, Powder Met., Vol 13, 1970, p 465-510 8. S. Oki, T. Akiyama, and K. Shoji, Fatigue Fracture Behavior of Sintered Carbon Steels, J. Japan Soc. Powder Met., Vol 30, 1983, p 229-234 9. W.B. James and R.C. O'Brien, High Performance Ferrous P/M Materials: The Effect of Alloying Method on Dynamic Properties, Progress in Powder Metallurgy, Vol 42, Metal Powder Industries Federation, 1986, p 353-372 10. H. Danninger, G. Jangg, B. Weiss, and R. Stickler, The Influence of Porosity on Static and Dynamic Properties of P/M Iron, PM into the 1990's, Vol 1, Proceedings of the World Conference on Powder Metallurgy, Institute of Materials, London, 1990, p 433-439 11. J. Holmes and R.A. Queeney, Fatigue Crack Initiation in a Porous Steel, Powder Met., Vol 28, 1985, p 231-235 12. S. Timoshenko and J.N. Goodier, Theory of Elasticity, McGraw-Hill, 1951, p 359-362 13. F.A. McClintock and A.S. Argon, Mechanical Behavior of Materials, Addison Wesley, 1966, p 412 14. C.M. Sonsino, F. Muller, V. Arnhold, and G. Schlieper, Influence of Mechanical Surface Treatments on the Fatigue Properties of Sintered Steels under Constant and Variable Stress Loading, Modern Developments in Powder Metallurgy, Vol 21, Metal Powder Industries Federation, 1988, p 55-66 15. C.M. Sonsino, G. Schlieper, and W.J. Huppmann, How to Improve the Fatigue Properties of Sintered Steels by Combined Mechanical and Thermal Treatments, Modern Developments in Powder Metallurgy, Vol 16, Met al Powder Industries Federation, 1985, p 33-48 16. J.H. Lange, M.F. Amateau, N. Sonti, and R.A. Queeney, Rolling Contact Fatigue in Ausrolled 1%C 9310 Steel, Inter. J. Fatigue, Vol 16, 1994, p 281-286 17. H. Kuroki and Y. Tokunaga, Effect of Density and Pore Shape on Impact Properties of Sintered Iron, Inter. J. Powder Met. Powder Tech., Vol 21, 1985, p 131-137 18. F.J. Esper and C.M. Sonsino, Fatigue Design for PM Components, European Powder Metallurgy Association, Shrewsbury, UK, 1994 19. K.D. Christian and R.M. German, Relation between Pore Structure and Fatigue Behavior in Sintered Iron- Copper-Carbon, Inter. J. Powder Met., Vol 31, 1995, p 51-61 20. I. Bertilsson, B. Karlsson, and J. Wasen, Fatigue Properties of Sintered Steels, Modern Developments in Powder Metallurgy, Vol 16, Metal Powder Industries Federation, 1985, p 19-32 21. R.C. O'Brien, Impact and Fatigue Characterization of Selected Ferrous P/M Materials, Progress in Powder Metallurgy, Vol 43, Metal Powder Industries Federation, 1987, p 749-775 22. P.S. Dasgupta and R.A. Queeney, Fatigue Crack Growth Rates in a Porous Metal, Inter. J. Fatigue, Vol 3, 1980, p 113-117 23. T. Prucher, Fatigue Life as a Function of the Mean Free Path between Inclusions, Modern Developments in Powder Metallurgy, Vol 18, Metal Powder Industries Federation, 1988, p 143-154 24. K.D. Christian, R.M. German, and A.S. Paulson, Statistical Analysis of Density and Particle Size Influences on Microstructural and Fatigue Properties of a Ferrous Alloy, Modern Developments in Powder Metallurgy, Vol 21, Metal Powder Industries Federation, 1988, p 23-39 25. I.J. Mellanby and J.R. Moon, The Fatigue Properties of Heat- Treatable Low Alloy Powder Metallurgy Steels, Modern Developments in Powder Metallurgy, Vol 18, Metal Powder Industries Federation, 1988, p 183-195 26. J.T. Barnby, D.C. Ghosh, and K. Dinsdale, Fracture Resistance of a Range of Steels, Powder Met., Vol 16, 1973, pp 55-71 27. E. Klar, D.F. Berry, P.K. Samal, J.J. Lewandowski, and J.D. Rigney, Fracture Toughness a nd Fatigue Crack Growth Response of Copper Infiltrated Steels, Inter. J. Powder Met., Vol 31, 1995, p 316-324 28. R.A. Queeney, Fatigue and Fracture Response of Metal-Infiltrated Sintered Powder Metals, Proceedings of ICM3, Vol 3, Pergamon Press, Oxford, 1979, p 373-381 29. D.A. Gerard and D.A. Koss, The Influence of Porosity on Short Fatigue Crack Growth at Large Strain Amplitudes, Inter. J. Fatigue, Vol 13, 1991, p 345-352 30. P.C. Paris, Fatigue An Interdisciplinary Approach, Syracuse University Press, 1964, p 107-117 31. W. Pompe, G. Leitner, K. Wetzig, G. Zies, and W. Grabner, Crack Propagation and Processes Near Crack Tip of Metallic Sintered Materials, Powder Met., Vol 27, 1984, p 45-51 Fatigue and Fracture Control for Powder Metallurgy Components * Randall M. German and Richard A. Queeney, The Pennsylvania State University Other Factors Determining Fatigue and Fracture Resistance It is well established that porosity is the major detriment to fatigue life for P/M materials. Beyond porosity, the sintered microstructure is a factor. Even in full-density materials fabricated by hot isostatic pressing, microstructure has a role. Weak links in the microstructure become evident during fracture. For porous structures, these weak links prove to be microstructural inhomogeneities, typically resulting from incomplete diffusional homogenization. Often powders (such as iron, nickel, and graphite) are mixed in the compaction stage. During heating the intent is for the mixed powders to homogenize to form a uniform microstructure, but this often is inhibited by too short a hold time at the peak temperature. Consequently, the alloying elements are poorly distributed and give point-to-point composition and microstructure changes, which are especially evident in postsintering heat treatment response. Accordingly, during fatigue or fracture, the weak links become the preferred failure paths. Most notable are the negative effects from inadequate homogenization of carbon to ensure uniform strength (Ref 18). Little is known about the sensitivity of fatigue and fracture to loading conditions for P/M material. Table 11 compares the 2 × 10 6 fatigue endurance strengths for Fe-1.5Cu-0.6C at 7.1 g/cm 3 density using bending and axial fatigue tests. The table also includes a comparison with two loading stress ratios (half-cycle and fully reversed, R = -1) and two notch conditions (unnotched and notched) (Ref 18). In these cases the unnotched loading shows little sensitivity to axial versus bending fatigue, but a large sensitivity is evident in the presence of notches. The notch sensitivity factor is reported to range between 0.32 and 0.43 for many of the common pressed and sintered P/M alloys (Ref 32). Table 11 Fatigue properties of Fe-1.5Cu-0.6C Notch factor, K Stress ratio, R Endurance strength, MPa at 2 × 10 6 cycles Axial -1 165 1.0 0 130 -1 84 2.8 0 64 Bending -1 160 1.0 0 127 -1 137 2.8 0 102 Note: Sintered to 7.2 g/cm 3 at 1120 °C for 30 min. Elastic modulus, 153 GPa; yield strength, 418 MPa; tensile strength, 483 MPa. Source: Full-density materials also suffer from residual microstructure artifacts that degrade the microstructure. In hot isostatically compacted powders, the achievement of 100% density is still insufficient to guarantee competitive fracture toughness, fatigue life, or even impact toughness. Thermally induced porosity is a subtle problem in many full-density P/M products. After consolidation the material is pore free, but it may contain small quantities of adsorbed gas. Once the product is put into high- temperature heat treatment or service, this residual gas precipitates to form pores if there is no compressive stress. In hot isostatically pressed titanium alloys, gas precipitation reportedly gives a 10 to 20% decrement in fatigue endurance strength (Ref 33). Another difficulty rests in slight contaminants located on the interfaces that were previously particle surfaces, a feature termed prior particle boundary decorations. Figure 2 shows such decorations in a fully densified P/M steel. Improper powder handling or cleaning prior to consolidation are the primary detriments. These contaminants remain on the powder interfaces, even though the structure is fully densified. Consequently, a small contamination film runs throughout the structure, providing an easy fracture path that is often traced to a trivial impurity level. The fracture path is along the prior particle boundaries and has a characteristic morphology, as shown in Fig. 3. In hot isostatically pressed Ti-6Al-4V there is substantial fatigue life improvement due to removal of the contaminant, with a change from 450 MPa endurance strength to 600 MPa due to powder cleaning prior to consolidation (Ref 34). Fig. 2 Prior particle boundary precipitates formed on a hot isostatically pressed st eel as the result of contamination during powder fabrication. 500× Fig. 3 A fracture surface showing preferential failure along prior particle boundaries. 150× One option for limiting the detrimental effects from prior particle boundary decorations is to forge the structure after hot isostatic pressing, a process often used in producing aerospace structures to ensure ultimate reliability. The forging operation upsets the microstructure and breaks apart the continuous films of contamination. The alternative is to resort to clean handling and processing, where the powder is produced by rapid solidification and kept under inert conditions during handling. These steps, which minimize segregation and contamination, are employed in the production of aerospace components, microelectronic structures, and high-performance filters. References cited in this section 15. C.M. Sonsino, G. Sch lieper, and W.J. Huppmann, How to Improve the Fatigue Properties of Sintered Steels by Combined Mechanical and Thermal Treatments, Modern Developments in Powder Metallurgy, Vol 16, Metal Powder Industries Federation, 1985, p 33-48 18. F.J. Esper and C.M. Sonsino, Fatigue Design for PM Components, European Powder Metallurgy Association, Shrewsbury, UK, 1994 32. A.F. Kravic, The Fatigue Properties of Sintered Iron and Steel, Inter. J. Powder Met., Vol 3 (No. 2), 1967, p 7- 13 33. R.L. Dreshfield and R.V. Miner, Effects of Thermally Induced Porosity on an as- HIP Powder Metallurgy Superalloy, Powder Met. Inter., Vol 12, 1980, p 83-87 34. F.H. Froes and C. Suryanarayana, Powder Processing of Titanium Alloys, Reviews in Particulate Materials, Vol 1, A. Bose, R.M. German, and A. Lawley, Ed., Metal Powder Industries Federation, 1993, p 223-276 Fatigue and Fracture Control for Powder Metallurgy Components * Randall M. German and Richard A. Queeney, The Pennsylvania State University Steps to Improve Fatigue and Fracture Resistance Surface pores are particularly detrimental to sintered materials with respect to fatigue life (Ref 35). Accordingly, carbonitriding and other surface strengthening and sealing treatments are most useful. Common treatments include shot peening, case hardening, repressing and resintering, coining, sizing, surface ausrolling, and postsintering heat treatments. For example, in pressed and sintered ferrous alloys, the endurance limit can be increased on small cross sections (6 by 6 mm) by at least 20% through shot peening. Carbonitriding is even more effective and can double the fatigue endurance limit. A typical carbonitride cycle involves heating to 940 °C in a mixture of ammonia and carbon dioxidefor 4 h to form a 0.5 mm deep carbon-rich layer. Surface grinding is another approach to improved fatigue strength. The larger the cross section of the material, the less benefit possible from surface treatments, because bulk material states will dominate mechanical response. The double press and double sinter approach was largely the only viable option for improving density and strength in traditional press and sinter P/M. This is more costly and involves extra tooling. A newly employed technique for improved fatigue life and fracture toughness in pressed and sintered ferrous alloys is to sinter at higher temperatures. The typical sintering temperature for steels is about 1120 °C, largely because of conveyor belt limitations in the furnaces. New materials of construction (ceramic belts) and new conveyor mechanisms (pusher plate and walking beam designs) allow higher- temperature processing regimes. Additionally, vacuum sintering usually is not limited in temperature, so it is viable for high- performance components. There is more sintering densification at the higher temperatures, so the density gain alone improves properties. Induced changes in the pore shape and size also improve fracture and fatigue properties. Several examples of the property gains are evident in Table 8. In a comparison of density gains versus sintering temperature effects, it is usually concluded that a change from 1120 to 1280 °C is equivalent to a density gain from 7.1 to 7.4 g/cm 3 in terms of both fracture toughness and fatigue endurance strength. For small components the surface treatments are most useful, because the compressive forces extend through a major portion of the microstructure. However, for the porous materials with large cross sections, the need is to sinter at higher temperatures to improve fatigue and fracture. Further, designs that minimize density gradients will assist in minimizing fatigue failure. The high-density regions have a higher fatigue strength, and the difference in strength with density often results in failure at the interface between high- and low-density regions. For fatigue-sensitive components, the tolerable range of densities is less than 0.05 g/cm 3 within the structure. As with all fatigue-sensitive components, consideration must be given to surface finishing and processing optimization. The keys to improved performance are reduction in the total porosity, elimination of segregation and contamination, and manipulation of the pore microstructure. [...]... 0.22 0.27 0.03 0.22 0.06 0.22 0.06 0.22 0.60 N 6 5-6 7 6 6-6 8 6 7-6 9 6 4-6 6 6 5-6 7 6 4-6 6 6 4-6 6 6 5-6 7 6 6-6 8 6 6-6 8 6 6-6 8 6 6-6 8 6 7-6 9 6 5-6 7 6 5-6 7 6 7-6 9 6 7-6 9 6 4-6 6 6 5-6 6 1.30 1.30 0.50 1.50 9.00 9.75 5.75 4.00 0.03 0 .07 5 3-5 5 6 0-6 2 5 7-5 9 5 9-6 3 4.25 1.30 0.40 1.05 2.10 4.25 4 2-4 8 4 4-5 2 HCHS, high carbon, high sulfur; HS, high sulfur Figure 8 schematically illustrates... German, and A.S Paulson, Statistical Analysis of Density and Particle Size Influences on Microstructural and Fatigue Properties of a Ferrous Alloy, Modern Developments in Powder Metallurgy, Vol 21, Metal Powder Industries Federation, 1988, p 2 3-3 9 25 I.J Mellanby and J.R Moon, The Fatigue Properties of Heat-Treatable Low Alloy Powder Metallurgy Steels, Modern Developments in Powder Metallurgy, Vol 18, Metal. .. Reviews in Particulate Materials, Vol 1, A Bose, R.M German, and A Lawley, Ed., Metal Powder Industries Federation, 1993, p 22 3-2 76 35 J.M Wheatley and G.C Smith, The Fatigue Strength of Sintered Iron, Powder Met., Vol 6, 1963, p 14 1-1 53 36 U Engstrom, C Lindberg, and I Tengzelius, Powders and Processes for High Performance PM Steels, Powder Met., Vol 35, 1992, p 6 7-7 2 Wear Resistance of Powder Metallurgy... under Constant and Variable Stress Loading, Modern Developments in Powder Metallurgy, Vol 21, Metal Powder Industries Federation, 1988, p 5 5-6 6 15 C.M Sonsino, G Schlieper, and W.J Huppmann, How to Improve the Fatigue Properties of Sintered Steels by Combined Mechanical and Thermal Treatments, Modern Developments in Powder Metallurgy, Vol 16, Metal Powder Industries Federation, 1985, p 3 3-4 8 16 J.H Lange,... bal bal bal bal bal Ni 10 bal Fe Cr 4 1 7-2 1 28 31 29 Other 2 4-2 6NiCr 1 4-1 6FeCrAlY Mo 2. 8-3 .3, Nb 4. 8-5 .5, Fe 1 4-2 1, Ti 0.7 5-1 .15 4W-1Si-1C 12.5W-1Si-2.5C 8.5W-1.5Si-1.5C Macrocrystalline tungsten carbide powders are a special kind of tungsten carbide powder manufactured by high temperature thermit process during which ore concentrate is converted directly... reference only For example, the P-M-K system found in Ref 5, and the "C" system in which C-1 through C4 are used for abrasion resistant grades and C-5 through C-8 are used for crater and deformation resistant grades C-1 through C-4 are straight WC-Co grades, and C-5 through C-8 contain cubic carbides, such as TiC and TaC for cutting tool applications A schematic flowchart of processes for cemented... Propagation and Processes Near Crack Tip of Metallic Sintered Materials, Powder Met., Vol 27, 1984, p 4 5-5 1 32 A.F Kravic, The Fatigue Properties of Sintered Iron and Steel, Inter J Powder Met., Vol 3 (No 2), 1967, p 713 33 R.L Dreshfield and R.V Miner, Effects of Thermally Induced Porosity on an as-HIP Powder Metallurgy Superalloy, Powder Met Inter., Vol 12, 1980, p 8 3-8 7 34 F.H Froes and C Suryanarayana, Powder. .. consolidating constituent powders together Particulate MMC is especially suitable for P/M processing In a narrower sense, MMC refers to modern low-density high-strength alloy systems such as aluminum- or titanium-base composites (e.g., Al-SiCp, Al-Al2O3, Ti-SiCp, and Ti-6Al-4V-SiCp) However, it is clear from the previous discussion that many P/M materials for wear resistant applications bear the same... Hardfacing and Thermal Spray Applications Metal alloy powders are used for hardfacing and thermal spray coatings for wear-resistant applications Hardfacing is the application of hard, wear-resistant material to the surface of a component by welding, thermal spraying, or a similar process for the main purpose of reducing wear Metal alloy powders used for hardfacing include WC-Co powders, other carbide -metal. .. Sintered Iron, Part 1: Basic Considerations and Review of Literature, Powder Met Inter., Vol 25, 1993, p 11 1-1 17 6 B Weiss, R Stickler, and H Sychra, High-Cycle Fatigue Behaviour of Iron-Base PM Materials, Metal Powder Report, Vol 45, 1990, p 18 7-1 92 7 R Haynes, Fatigue Behaviour of Sintered Metals and Alloys, Powder Met., Vol 13, 1970, p 46 5-5 10 8 S Oki, T Akiyama, and K Shoji, Fatigue Fracture Behavior . 148 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1120 °C, 2 h 7.1 Rotating R = -1 135 Fe-4Ni-1.5Cu-0.5Mo-0.6C P+S, 1250 °C, h 7.1 Rotating R = -1 147 Rotating R = -1 195 Fe-4Ni-1.5Cu-0.5Mo-0.6C. toughness, MPa Fe-4.4Cr-9.2Co-7.2V-3.7Mo-9.2W-2.7C P+S, 1150 °C, 1 h 8.1 13 Fe-1.5Cu-2Ni-0.8C P+S, 1120 °C, h 6.8 40 Fe-1.8Ni-0.5Mo-1.5C P+S, 1120 °C, h 6.6 15 Fe-1.8Ni-0.5Mo-1.5C P+S,. Fe-1.8Ni-0.5Mo-1.5C P+S, 1120 °C, h 7.1 24 Fe-1.8Ni-0.5Mo-1.5C DP + DS, 1100 °C, h 7.5 2 1-3 8 Fe-1.8Ni-0.5Mo-1.5C HF, 1100 °C 7.85 64 Fe-0.8P-0.3C P+S, 1120 °C, h 7.0 22 Fe-0.8P-0.3C