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Dynamic embrittlement–time-dependent brittle fracture
Journal of Alloys and Compounds 378 (2004) 79–84
Dynamic embrittlement–time-dependent brittle fracture U. Krupp a,∗ , C.J. McMahon, Jr b b
a Institut für Werkstofftechnik, Universität Siegen, 57068 Siegen, Germany Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia 19104, USA
Received 3 September 2003; accepted 14 October 2003
Abstract Dynamic embrittlement (DE) is a generic type of brittle intergranular fracture where the crack propagation is controlled by grain boundary diffusion of an embrittling element, which can come from the material itself, e.g., during S-induced stress relief cracking of steels, or from the surrounding atmosphere, e.g., hold-time cracking of Ni-base superalloys. Four-point bend tests on poly- and bicrystalline specimens revealed that cracking by dynamic embrittlement depends strongly on local microstructural features, e.g., the grain-boundary structure. © 2004 Elsevier B.V. All rights reserved. Keywords: Metals; Grain boundaries; Brittle fracture; Dynamic embrittlement
1. Introduction When we speak of brittle fracture of metallic materials, we normally think of either transcrystalline cleavage or intergranular decohesion, both of which involve crack propagation rates on the order of 103 m/s. Plasticity is usually involved in the nucleation of the crack, but during propagation any plasticity that occurs is merely coincidental. Dynamic embrittlement (DE) is another mode of brittle fracture that usually takes place by decohesion along grain boundaries, but here the propagation rates are typically on the order of 10−7 to 10−5 m/s. Again, plastic flow during propagation is only coincidental. The process of crack extension involves the inward diffusion of a mobile embrittling element from a free surface as a result of the application of a tensile stress. It is essentially analogous to the phenomenon of diffusion-controlled creep-cavity growth a la Hull and Rimmer [1]. As illustrated by Fig. 1, the difference is that in DE it is the surface-adsorbed embrittling element that diffuses into the grain boundary, rather than matrix atoms. The driving force for the diffusion in both cases is the same: the reduction in chemical potential of an atom when it enters the grain boundary from the surface (the work done by the
∗ Corresponding author. Tel.: +49-271-740-2184; fax: +49-271-740-2545. E-mail address: [email protected] (U. Krupp).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.10.074
applied stress σ is equal to σΩ, Ω being the atomic volume of the diffusing atom).
2. Two categories of dynamic embrittlement In one category of DE the embrittling element comes from the material itself, by way of segregation to the surface. Examples are sulfur-induced cracking in alloy steels and tin-induced cracking in Cu–Sn alloys [2–5]. In both cases, the cracking appeared to initiate internally around cavities formed at intergranular inclusions during plastic flow, as illustrated schematically in Fig. 2. In the case of the steel, the sulfur content in the grain-boundary regions has been enhanced by the thermal history (see below), and there is some circumstantial evidence for tin enrichment of the grain boundaries in the Cu–Sn alloy [6]. In the other type of DE the embrittling element comes from the environment. The most-studied case so far is oxygen-induced cracking of nickel-alloys [7–13], although the same thing has been observed in Cu–0.25%Be [14]. The term “dynamic embrittlement” was first used to describe oxygen-induced cracking in Ni3 Al [12]. Most other studies have been done using cyclic loading, and the general finding has been that it involves time-dependent, rather than cycle-dependent, crack propagation [7]. The rate of cracking increases as the cyclic frequency decreases, as the temperature increases, and as the oxygen pressure increases. The amount of cracking is enhanced by hold-times under
80
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
stress is two-fold: first, to induce the penetration of the solid by the surface-adsorbed embrittling element, and, second, to cause the bond stretching and decohesion in the embrittled region ahead of the crack or void.
3. Cracking from internal elements
Fig. 1. (a) Schematic representation of Hull–Rimmer cavity growth. (b) Schematic illustration of the stress-induced grain-boundary penetration that is characteristic of dynamic embrittlement.
tensile stress. In fact, in the gas-turbine industry the phenomenon is known as “hold-time” cracking. There is reason to believe that liquid-metal embrittlement is an extreme form of dynamic embrittlement, in which the concentration and mobility of the embrittling element is extremely high [15]. An analogous phenomenon would be hot tearing in castings. Cracking from coatings of low-melting metals just below their freezing point has been called “solid-metal embrittlement” to contrast the much faster cracking with that just above the freezing point [16], and this all has the appearance of generic dynamic embrittlement. In addition to the surface concentration of the embrittling element, the important factors are the level of tensile stress, the mobility of the embrittling element (as determined by the temperature), and the grain-boundary diffusivity of the embrittling element in the material in question. The role of
The prototypical case is that of stress-relief cracking of alloy steels, which tends to occur in the heat-affected zones of weldments while the structure is being heated for relief of internal stresses. The sulfur is made available to the grain boundaries by the high-temperature excursion during the welding, in which sulfides dissolve and sulfur segregates to the boundaries. As the grain-coarsened HAZ is cooled by heat flow to the surrounding metal, the sulfur either remains as segregated atoms or forms small (metastable) sulfides along the boundaries. The creep that occurs during stress relief can open cavities around inclusions. Sulfur spreads quickly along the surfaces of the cavities, and, if the stress is high enough, brittle cracking can initiate from the cavities [2] (cf. Fig. 2a). If the stress is not high enough, much slower cracking occurs by rupture along the array of fine sulfides lining the grain boundaries [2]. A later study of DE [4] has shown that the cracking occurs by percolation of sharp cracks along boundaries in which the embrittling element has a high diffusivity. Initially, un-cracked ligaments are left behind the main crack front, presumably because no fast-diffusion boundaries were available in these locations. The crack velocity v versus the stress intensity factor K curve shows bursts of cracking initially, presumably along patches of fast-diffusion boundaries, as shown in Fig. 3 or in Fig. 4a (da/dt versus time). The locations of the un-cracked ligaments in an interrupted specimen are shown in Fig. 4b. The influence of grain-boundary structure on the rate of cracking was studied in a Cu–8%Sn alloy, which shows a behavior analogous to the alloy steel. In this case,
Fig. 2. Schematic illustration of nucleation of DE at cavities around inclusions in (a) steel and (b) a Cu–Sn alloy.
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
81
Fig. 3. Example of crack growth in bursts in the early stages of sulfur-induced dynamic embrittlement (stress-relief cracking) in a MnNiCrMo steel at 560 ◦ C, ref. [4].
surface-segregated tin is the embrittling element [3]. Bicrystals having a symmetrical 5 [0 0 1] (0 3 1) tilt boundary were fabricated from this alloy by diffusion bonding of single crystals, and the cracking of the boundary was studied
in the fast-diffusion direction parallel to the tilt axis and in the slow-diffusion direction perpendicular to the tilt axis [17]. It was found that the bicrystals cracked readily along the tilt axis, as long as there was a sufficient number of Al2 O3 particles in the boundary to nucleate voids. However, brittle cracking did not occur at all perpendicular to the tilt axis. In the former case, the cracking was continuous, and it occurred at stress intensities of less than 3 MPa m0.5 .
4. Cracking from environmental elements
Fig. 4. Percolation-type intergranular fracture: (a) discontinuous crack propagation and (b) location of un-cracked ligaments left behind the main crack in the fracture surface after an interrupted test. The specimen of the same steel as in Fig. 3 was broken open by impact, ref. [4].
Time-dependent cracking during cyclic loading of the commonly used superalloy IN718 has been studied by a number of authors [7–11] and is generally attributed to oxidation effects. More recently it has been shown in static-loading experiments [13] that this kind of cracking seems to be the same phenomenon as dynamic embrittlement. Because of their relative simplicity, the static-loading experiments are easier to interpret and so will be our focus here. The cracking of notched specimens loaded in pure bending under a fixed displacement can be monitored by the drop in the load as a function of time, as shown in Fig. 5a. This can be converted to crack length versus time by means of a compliance calibration of the specimen, and consequently to crack velocity v = da/dt versus the stress intensity factor as shown in Fig. 5(b). The effect of oxygen pressure is apparent in these figures, but it can be shown more dramatically by an experiment in which the oxygen is suddenly removed from the test chamber and later suddenly re-admitted, as
82
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
Fig. 6. Part of a load-relaxation curve for alloy 718 at 650 ◦ C showing the immediate response to the removal and re-introduction of oxygen in the test chamber, ref. [13].
Fig. 5. (a) Load-relaxation curves due to intergranular cracking in IN718 at 650 ◦ C at various pressures of oxygen. (b) da/dt versus K curves derived from curves as in (a), ref. [13].
shown in Fig. 6. From this it is clear that the zone of action of the oxygen must be very close to the crack tip. When a steady-state model [18] is applied to the governing differential equation for DE, i.e., stress-assisted GB diffusion into the cohesive zone in front of an intercrystalline crack tip, the latter being the origin of the x-coordinate ∂c ∂2 c DGB ∂ ∂σ = DGB 2 − c , (1) ∂t kT ∂x ∂x ∂x an approximate solution can be found for the embrittler concentration c ahead of the crack tip v DGB c = c0 exp − x+ (2) (σ − σ0 ) . DGB kT Assuming the oxygen GB diffusivity to be in the range between DGB = 10−14 and 10−15 m2 /s, σ−σ 0 to be close to the yield strength, the measured crack velocity of v = 30 m/s
can be explained only if the extent of oxygen penetration by grain boundary diffusion is on the nanometer scale [13,19]. As in the case of steel discussed above, the rate of cracking varies from one grain boundary to another, depending on the ease of diffusion; that is, depending on the grain-boundary structure. Because of that, cracking in polycrystalline IN718 occurs discontinuously, as shown in Fig. 7. As a crack advances in one location, the local stress intensity drops, and the load is shifted to other regions that are not cracking. As creep occurs in these regions, the stress intensity again increases, and another increment of cracking occurs. The fact, that the experimentally-determined temperature dependence of the cracking rate matches approximately that of self diffusion in nickel confirms the plausibility of the mechanism proposed above that includes power-law creep [13]. In addition, the applied stress intensity of tens of MPa m0.5 correlates with the creep process, rather than with dynamic embrittlement. This is inherent in experiments on polycrystals; experiments on bicrystals are necessary if one wants to study the effects of temperature and stress intensity on the cracking process itself. Preliminary experiments have been done on bicrystals of IN718 [19], and it has been found that a random (i.e., fast-diffusion) boundary cracks much more rapidly than a symmetrical 5 (0 3 1) [0 0 1] tilt boundary, as shown in Fig. 8. The rapid cracking of the random boundary appears to have been continuous, without the arrest marks typically found in polycrystals. In view of the effect of grain-boundary structure, it might be expected that thermomechanical processing (TMP) [20–22] aimed at increasing the fraction of special boundaries that resist diffusion should improve the resistance to dynamic embrittlement. This has now been shown to be the case in IN718 [23], as seen in Fig. 9. Here, four cycles of cold rolling and annealing have increased the fraction of 3
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
83
Fig. 7. The tip region of a crack that was growing in alloy 718 in 1 atm oxygen at 650 ◦ C, the growth of which was interrupted by unloading the specimen. The lower part of the micrograph represents post-test impact fracture at room temperature, ref. [13].
boundaries from about 16 to 34%, and this has obviously provided increased resistance to dynamic embrittlement as compared to the as-received condition (AR).
5. Summary Dynamic embrittlement appears to be a generic form of brittle fracture. It’s unique feature is that it occurs quasi-statically, so its kinetics can be readily studied. It occurs when a mobile surface-adsorbed embrittling element
Fig. 8. Comparison of the load-relaxation curves due to cracking in a bicrystal with a random boundary vs. a symmetrical 5 tilt boundary, ref. [19].
is present on the surface of a material to which a sufficient tensile stress is applied. The embrittling element can originate from within the alloy itself and get to the surface by segregation, or it can come from the environment. Cracking by a liquid metal may be an extreme form of dynamic embrittlement [15] in which the concentration and mobility of the embrittling element are extraordinarily high. However, this aspect needs more study.
Fig. 9. Effect of grain-boundary-engineering-type thermomechanical processing (TMP) on the cracking behavior of IN718 in air at 650 ◦ C, ref. [23].
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Acknowledgements This research has been supported by the U.S. Department of Energy, Basic Energy Sciences, under grant no. DE-FG02-01ER45924, the U.S. Air Force Office of Scientific Research under grant no. 538532, and by the Alexander von Humboldt Foundation through a Feodor Lynen fellowship to U. Krupp.
References [1] D. Hull, D.E. Rimmer, Phil. Mag. 4 (1959) 673. [2] J.S. Shin, C.J. McMahon Jr., Acta Metall. 32 (1984) 1535. [3] E.V. Barrera, M. Menyhard, D. Bika, B. Rothman, C.J. McMahon Jr., Scripta Metall. 27 (1992) 205. [4] D. Bika, J.A. Pfaendtner, M. Menyhard, C.J. McMahon Jr., Acta Metall. et Mater. 43 (1995) 1909. [5] D. Bika, C.J. McMahon Jr., Acta Metall. et Mater. 43 (1995) 1895. [6] J. Lerner, S.S. Woods, E. Peng, C.J. McMahon Jr., Mater. Sci. Eng. A345 (2003) 357.
[7] K.-M. Chang, M.F. Henry, M.G. Benz, JOM 12 (1990) 29. [8] H. Ghonem, T. Nicholas, A. Pineau, Fatigue Fract. Eng. Mater. Struct. 16 (1993) 565. [9] H. Ghonem, T. Nicholas, A. Pineau, Fatigue Fract. Eng. Mater. Struct. 16 (1993) 577. [10] R. Molins, G. Hochstetter, J.C. Chassaigne, E. Andrieu, Acta Mater. 45 (1997) 663. [11] H. Ghonem, D. Zheng, Mater. Sci. Eng. A150 (1992) 151. [12] C.T. Liu, C.L. White, Acta Metall. 35 (1987) 643. [13] J.A. Pfaendtner, C.J. McMahon Jr., Acta Mater. 49 (2001) 3369. [14] R.C. Muthiah, C.J. McMahon Jr., A. Guha, Mater. Sci. Forum. 207–209 (1996) 585. [15] J. Lerner, C.J. McMahon Jr., Mater. Sci Eng. A336 (2002) 72. [16] P. Gordon, H.H. An, Metall. Trans. A 13A (1982) 457. [17] R.C. Muthiah, J.A. Pfaendtner, S. Ishikawa, C.J. McMahon Jr., Acta Mater. 47 (1999) 2797. [18] Y. Xu, J.L. Bassani, Mater. Sci. Eng. A260 (1999) 48. [19] U. Krupp, Int. Mater. Rev., in press. [20] G. Palumbo, E.M. Lehockey, P. Lin, JOM 50 (2) (1998) 40. [21] G.S. Was, V. Thaveeprungsriporn, D.C. Crawford, JOM 50 (2) (1998) 44. [22] M. Kumar, W.E. King, A.J. Schwartz, Acta Mater. 48 (2000) 2081. [23] U. Krupp, W.M. Kane, X-Y. Liu, O. Dueber, C. Laird, C.J. McMahon, Mater Sci. Eng. A349 (2003) 213.
Dynamic embrittlement–time-dependent brittle fracture U. Krupp a,∗ , C.J. McMahon, Jr b b
a Institut für Werkstofftechnik, Universität Siegen, 57068 Siegen, Germany Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia 19104, USA
Received 3 September 2003; accepted 14 October 2003
Abstract Dynamic embrittlement (DE) is a generic type of brittle intergranular fracture where the crack propagation is controlled by grain boundary diffusion of an embrittling element, which can come from the material itself, e.g., during S-induced stress relief cracking of steels, or from the surrounding atmosphere, e.g., hold-time cracking of Ni-base superalloys. Four-point bend tests on poly- and bicrystalline specimens revealed that cracking by dynamic embrittlement depends strongly on local microstructural features, e.g., the grain-boundary structure. © 2004 Elsevier B.V. All rights reserved. Keywords: Metals; Grain boundaries; Brittle fracture; Dynamic embrittlement
1. Introduction When we speak of brittle fracture of metallic materials, we normally think of either transcrystalline cleavage or intergranular decohesion, both of which involve crack propagation rates on the order of 103 m/s. Plasticity is usually involved in the nucleation of the crack, but during propagation any plasticity that occurs is merely coincidental. Dynamic embrittlement (DE) is another mode of brittle fracture that usually takes place by decohesion along grain boundaries, but here the propagation rates are typically on the order of 10−7 to 10−5 m/s. Again, plastic flow during propagation is only coincidental. The process of crack extension involves the inward diffusion of a mobile embrittling element from a free surface as a result of the application of a tensile stress. It is essentially analogous to the phenomenon of diffusion-controlled creep-cavity growth a la Hull and Rimmer [1]. As illustrated by Fig. 1, the difference is that in DE it is the surface-adsorbed embrittling element that diffuses into the grain boundary, rather than matrix atoms. The driving force for the diffusion in both cases is the same: the reduction in chemical potential of an atom when it enters the grain boundary from the surface (the work done by the
∗ Corresponding author. Tel.: +49-271-740-2184; fax: +49-271-740-2545. E-mail address: [email protected] (U. Krupp).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.10.074
applied stress σ is equal to σΩ, Ω being the atomic volume of the diffusing atom).
2. Two categories of dynamic embrittlement In one category of DE the embrittling element comes from the material itself, by way of segregation to the surface. Examples are sulfur-induced cracking in alloy steels and tin-induced cracking in Cu–Sn alloys [2–5]. In both cases, the cracking appeared to initiate internally around cavities formed at intergranular inclusions during plastic flow, as illustrated schematically in Fig. 2. In the case of the steel, the sulfur content in the grain-boundary regions has been enhanced by the thermal history (see below), and there is some circumstantial evidence for tin enrichment of the grain boundaries in the Cu–Sn alloy [6]. In the other type of DE the embrittling element comes from the environment. The most-studied case so far is oxygen-induced cracking of nickel-alloys [7–13], although the same thing has been observed in Cu–0.25%Be [14]. The term “dynamic embrittlement” was first used to describe oxygen-induced cracking in Ni3 Al [12]. Most other studies have been done using cyclic loading, and the general finding has been that it involves time-dependent, rather than cycle-dependent, crack propagation [7]. The rate of cracking increases as the cyclic frequency decreases, as the temperature increases, and as the oxygen pressure increases. The amount of cracking is enhanced by hold-times under
80
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stress is two-fold: first, to induce the penetration of the solid by the surface-adsorbed embrittling element, and, second, to cause the bond stretching and decohesion in the embrittled region ahead of the crack or void.
3. Cracking from internal elements
Fig. 1. (a) Schematic representation of Hull–Rimmer cavity growth. (b) Schematic illustration of the stress-induced grain-boundary penetration that is characteristic of dynamic embrittlement.
tensile stress. In fact, in the gas-turbine industry the phenomenon is known as “hold-time” cracking. There is reason to believe that liquid-metal embrittlement is an extreme form of dynamic embrittlement, in which the concentration and mobility of the embrittling element is extremely high [15]. An analogous phenomenon would be hot tearing in castings. Cracking from coatings of low-melting metals just below their freezing point has been called “solid-metal embrittlement” to contrast the much faster cracking with that just above the freezing point [16], and this all has the appearance of generic dynamic embrittlement. In addition to the surface concentration of the embrittling element, the important factors are the level of tensile stress, the mobility of the embrittling element (as determined by the temperature), and the grain-boundary diffusivity of the embrittling element in the material in question. The role of
The prototypical case is that of stress-relief cracking of alloy steels, which tends to occur in the heat-affected zones of weldments while the structure is being heated for relief of internal stresses. The sulfur is made available to the grain boundaries by the high-temperature excursion during the welding, in which sulfides dissolve and sulfur segregates to the boundaries. As the grain-coarsened HAZ is cooled by heat flow to the surrounding metal, the sulfur either remains as segregated atoms or forms small (metastable) sulfides along the boundaries. The creep that occurs during stress relief can open cavities around inclusions. Sulfur spreads quickly along the surfaces of the cavities, and, if the stress is high enough, brittle cracking can initiate from the cavities [2] (cf. Fig. 2a). If the stress is not high enough, much slower cracking occurs by rupture along the array of fine sulfides lining the grain boundaries [2]. A later study of DE [4] has shown that the cracking occurs by percolation of sharp cracks along boundaries in which the embrittling element has a high diffusivity. Initially, un-cracked ligaments are left behind the main crack front, presumably because no fast-diffusion boundaries were available in these locations. The crack velocity v versus the stress intensity factor K curve shows bursts of cracking initially, presumably along patches of fast-diffusion boundaries, as shown in Fig. 3 or in Fig. 4a (da/dt versus time). The locations of the un-cracked ligaments in an interrupted specimen are shown in Fig. 4b. The influence of grain-boundary structure on the rate of cracking was studied in a Cu–8%Sn alloy, which shows a behavior analogous to the alloy steel. In this case,
Fig. 2. Schematic illustration of nucleation of DE at cavities around inclusions in (a) steel and (b) a Cu–Sn alloy.
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81
Fig. 3. Example of crack growth in bursts in the early stages of sulfur-induced dynamic embrittlement (stress-relief cracking) in a MnNiCrMo steel at 560 ◦ C, ref. [4].
surface-segregated tin is the embrittling element [3]. Bicrystals having a symmetrical 5 [0 0 1] (0 3 1) tilt boundary were fabricated from this alloy by diffusion bonding of single crystals, and the cracking of the boundary was studied
in the fast-diffusion direction parallel to the tilt axis and in the slow-diffusion direction perpendicular to the tilt axis [17]. It was found that the bicrystals cracked readily along the tilt axis, as long as there was a sufficient number of Al2 O3 particles in the boundary to nucleate voids. However, brittle cracking did not occur at all perpendicular to the tilt axis. In the former case, the cracking was continuous, and it occurred at stress intensities of less than 3 MPa m0.5 .
4. Cracking from environmental elements
Fig. 4. Percolation-type intergranular fracture: (a) discontinuous crack propagation and (b) location of un-cracked ligaments left behind the main crack in the fracture surface after an interrupted test. The specimen of the same steel as in Fig. 3 was broken open by impact, ref. [4].
Time-dependent cracking during cyclic loading of the commonly used superalloy IN718 has been studied by a number of authors [7–11] and is generally attributed to oxidation effects. More recently it has been shown in static-loading experiments [13] that this kind of cracking seems to be the same phenomenon as dynamic embrittlement. Because of their relative simplicity, the static-loading experiments are easier to interpret and so will be our focus here. The cracking of notched specimens loaded in pure bending under a fixed displacement can be monitored by the drop in the load as a function of time, as shown in Fig. 5a. This can be converted to crack length versus time by means of a compliance calibration of the specimen, and consequently to crack velocity v = da/dt versus the stress intensity factor as shown in Fig. 5(b). The effect of oxygen pressure is apparent in these figures, but it can be shown more dramatically by an experiment in which the oxygen is suddenly removed from the test chamber and later suddenly re-admitted, as
82
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
Fig. 6. Part of a load-relaxation curve for alloy 718 at 650 ◦ C showing the immediate response to the removal and re-introduction of oxygen in the test chamber, ref. [13].
Fig. 5. (a) Load-relaxation curves due to intergranular cracking in IN718 at 650 ◦ C at various pressures of oxygen. (b) da/dt versus K curves derived from curves as in (a), ref. [13].
shown in Fig. 6. From this it is clear that the zone of action of the oxygen must be very close to the crack tip. When a steady-state model [18] is applied to the governing differential equation for DE, i.e., stress-assisted GB diffusion into the cohesive zone in front of an intercrystalline crack tip, the latter being the origin of the x-coordinate ∂c ∂2 c DGB ∂ ∂σ = DGB 2 − c , (1) ∂t kT ∂x ∂x ∂x an approximate solution can be found for the embrittler concentration c ahead of the crack tip v DGB c = c0 exp − x+ (2) (σ − σ0 ) . DGB kT Assuming the oxygen GB diffusivity to be in the range between DGB = 10−14 and 10−15 m2 /s, σ−σ 0 to be close to the yield strength, the measured crack velocity of v = 30 m/s
can be explained only if the extent of oxygen penetration by grain boundary diffusion is on the nanometer scale [13,19]. As in the case of steel discussed above, the rate of cracking varies from one grain boundary to another, depending on the ease of diffusion; that is, depending on the grain-boundary structure. Because of that, cracking in polycrystalline IN718 occurs discontinuously, as shown in Fig. 7. As a crack advances in one location, the local stress intensity drops, and the load is shifted to other regions that are not cracking. As creep occurs in these regions, the stress intensity again increases, and another increment of cracking occurs. The fact, that the experimentally-determined temperature dependence of the cracking rate matches approximately that of self diffusion in nickel confirms the plausibility of the mechanism proposed above that includes power-law creep [13]. In addition, the applied stress intensity of tens of MPa m0.5 correlates with the creep process, rather than with dynamic embrittlement. This is inherent in experiments on polycrystals; experiments on bicrystals are necessary if one wants to study the effects of temperature and stress intensity on the cracking process itself. Preliminary experiments have been done on bicrystals of IN718 [19], and it has been found that a random (i.e., fast-diffusion) boundary cracks much more rapidly than a symmetrical 5 (0 3 1) [0 0 1] tilt boundary, as shown in Fig. 8. The rapid cracking of the random boundary appears to have been continuous, without the arrest marks typically found in polycrystals. In view of the effect of grain-boundary structure, it might be expected that thermomechanical processing (TMP) [20–22] aimed at increasing the fraction of special boundaries that resist diffusion should improve the resistance to dynamic embrittlement. This has now been shown to be the case in IN718 [23], as seen in Fig. 9. Here, four cycles of cold rolling and annealing have increased the fraction of 3
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
83
Fig. 7. The tip region of a crack that was growing in alloy 718 in 1 atm oxygen at 650 ◦ C, the growth of which was interrupted by unloading the specimen. The lower part of the micrograph represents post-test impact fracture at room temperature, ref. [13].
boundaries from about 16 to 34%, and this has obviously provided increased resistance to dynamic embrittlement as compared to the as-received condition (AR).
5. Summary Dynamic embrittlement appears to be a generic form of brittle fracture. It’s unique feature is that it occurs quasi-statically, so its kinetics can be readily studied. It occurs when a mobile surface-adsorbed embrittling element
Fig. 8. Comparison of the load-relaxation curves due to cracking in a bicrystal with a random boundary vs. a symmetrical 5 tilt boundary, ref. [19].
is present on the surface of a material to which a sufficient tensile stress is applied. The embrittling element can originate from within the alloy itself and get to the surface by segregation, or it can come from the environment. Cracking by a liquid metal may be an extreme form of dynamic embrittlement [15] in which the concentration and mobility of the embrittling element are extraordinarily high. However, this aspect needs more study.
Fig. 9. Effect of grain-boundary-engineering-type thermomechanical processing (TMP) on the cracking behavior of IN718 in air at 650 ◦ C, ref. [23].
84
U. Krupp, C.J. McMahon, Jr / Journal of Alloys and Compounds 378 (2004) 79–84
Acknowledgements This research has been supported by the U.S. Department of Energy, Basic Energy Sciences, under grant no. DE-FG02-01ER45924, the U.S. Air Force Office of Scientific Research under grant no. 538532, and by the Alexander von Humboldt Foundation through a Feodor Lynen fellowship to U. Krupp.
References [1] D. Hull, D.E. Rimmer, Phil. Mag. 4 (1959) 673. [2] J.S. Shin, C.J. McMahon Jr., Acta Metall. 32 (1984) 1535. [3] E.V. Barrera, M. Menyhard, D. Bika, B. Rothman, C.J. McMahon Jr., Scripta Metall. 27 (1992) 205. [4] D. Bika, J.A. Pfaendtner, M. Menyhard, C.J. McMahon Jr., Acta Metall. et Mater. 43 (1995) 1909. [5] D. Bika, C.J. McMahon Jr., Acta Metall. et Mater. 43 (1995) 1895. [6] J. Lerner, S.S. Woods, E. Peng, C.J. McMahon Jr., Mater. Sci. Eng. A345 (2003) 357.
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