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Deformation substructures and fracture behavior of CuSi alloys subjected to dynamic loading conditions
Materials Science and Engineering, 46 ( 1 9 8 0 ) 15 - 22 © Elsevier S e q u o i a S.A., L a u s a n n e - - P r i n t e d in t h e N e t h e r l a n d s
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Deformation Substructures and Fracture B e h a v i o r o f C u - S i A l l o y s S u b j e c t e d to Dynamic Loading C o n d i t i o n s K. V. R A O and D. H. P O L O N I S
Metallurgical Engineering Division, University of Washington, Seattle, Wash. 98195 (U.S,A.) (Received D e c e m b e r 3, 1 9 7 9 )
SUMMARY
The microstructures and fracture surfaces of dynamically loaded Cu-Si alloys were examined by means of transmission electron microscopy and scanning electron microscopy techniques. The features of the deformation and fracture processes are interpreted in terms of the microstructural damage resulting from dynamic loading. The occurrence of a straininduced phase transformation in dynamically loaded solution-treated alloys was confirmed by the results o f both transmission electron microscopy and X-ray diffraction analysis. The deformation substructure after dynamic loading exhibits a profusion of microtwins in the matrix and the preferred nucleation of microvoids along ~-~ interfaces. A ductile fracture mode characterizes the failure under both static and dynamic loading conditions, but a lesser ductile mode associated with smaller dimple sizes was evident in dynamically loaded alloys that had been aged to produce the ~ phase.
1. I N T R O D U C T I O N
The response of a given material to dynamic loading conditions depends on a combination of factors including the stress history, the microstructure, the dislocation mobility and the presence of discontinuities and flaws. The prediction of dynamic failure in materials subjected to high intensity stress waves is limited by an inability to understand h o w the fracture process is controlled by changes in the microstructure and constitution of a material. Uncertainty still exists regarding the particular fracture mechanism
or combination of mechanisms that apply when the microstructure of a metal is modified by heat treatment or when the stress field conditions are changed. The complexity of the problem is illustrated by the processes of void formation and coalescence which are often important steps in the ductile fracture mechanism under both static and dynamic loading conditions. Previous workers have shown that impurity second-phase particles usually control the nucleaticn of voids leading to spallation under dynamic loading conditions [1 - 4]. In order to resolve the influence of precipitation processes on the dynamic behavior of materials, it is necessary to minimize the influence of impurity second-phase particles by employing alloys of high purity and carefully controlled composition. In the present work the response of Cu-Si alloys to dynamic loading conditions was investigated. The addition of silicon reduces the stacking fault energy of copper and promotes the formation of planar dislocation arrays and expanded faults [ 5]. Compositions with more than 4.6% Si (all alloy compositions are expressed in weight percentages) exhibit the formation of the h.c.p. K phase after appropriate thermal or mechanical processing [6]. The ~ phase forms as platelets parallel to {lll}~.c.c.: these platelets act as barriers for cross-slip processes and are expected to influence the flow and fracture behavior accompanying dynamic loading. In this paper we report on the microstructural changes and fracture modes due to the dynamic loading of Cu-Si alloys after solution treatment and after aging to produce a twophase mixture of the a and K phases. The influence of ~ platelets on the microstructural origins of the voids leading to ductile fracture was of special interest in this work.
16 2. EXPERIMENTAL PROCEDURE
Cu-Si alloys containing 4.37% Si and 4.67% Si were prepared by vacuum melting and casting from OFHC copper and 99.999% Si. The cylindrical rods were homogenized at 750 °C for 48 h and subsequently cold swaged to rods 1.3 cm in diameter with intermediate annealing treatments. Hollow cylindrical test specimens with an inside diameter of 5 mm, an outside diameter of 10 mm and a length of 19.1 mm were machined from the homogenized rods and solution treated at 700 °C for 48 h. Aging treatments were conducted at 470 and 350 °C on the solution-treated and quenched specimens. Prior to heat treatment the specimens were encapsulated in Vycor tubing, evacuated and backfilled with a partial pressure of argon. An exploding wire system developed by Fyfe and Ensminger [7] was employed to generate the dynamic loading conditions in the hollow cylindrical specimens. The system consists of two 15 p F capacitors charged to a nominal 20 kV and providing an energy storage of 3000 J per capacitor for a total of 6000 J. Specimens loaded in this fashion expand in the radial direction with complete symmetry and with little variation along the axis of the cylinder. This loading technique results in a biaxial strain condition in thickwalled cylinders where the hoop and radial strains are non-zero. The high intensity stress waves required for the dynamic loading of thick-walled cylinders are generated by discharging the capacitors through a polyethylene-coated wire. The polyethylene is driven against the inside diameter of the specimen and acts as a cylindrical piston. The pressure pulse has a rise time of approximately 0.1 ps, and the maximum load can be as high as 25 kbar. The present experiments employed copper wires 0.84 mm in diameter encased in a polyethylene jacket 2.95 mm in diameter. The static tensile tests employed an Instron testing machine operated at a strain rate of 3.3 × 10 -4 s-1 at room temperature. The X-ray diffraction analyses were performed with a Philips vertical diffractometer and Cu Ks radiation generated at 40 kW and 20 mA. A graphite single-crystal m o n o c h r o m a t o r was used to discriminate the Cu Ks radiation. Thin foil specimens for transmission electron microscopy were prepared from slices
10 × 10 -3 in thick cut from the inner surface of the cylinders in a direction tangential to the circumference. The slices were chemically polished in a solution of 50% nitric acid, 25% acetic acid and 25% orthophosphoric acid. Discs 3.05 mm in diameter punched from these slices were electropolished in a jet polisher using a solution containing one part of nitric acid and two parts of methyl alcohol at temperatures below --20 °C. The thin foils were examined in a JEM-7 electron microscope operated at 100 kV. The examination of fracture surfaces was conducted by scanning electron microscopy (SEM) techniques.
3. RESULTS AND DISCUSSION
The 4.37% Si and 4.67% Si alloys used in the present study exhibited less macroscopic strain and a reduced Susceptibility to spallation compared with alloys of lower silicon content examined in an earlier study [8]. The Cu-4.37%Si alloy is close to the limit above which the supersaturated ~ phase in solutiontreated alloys is partially transformed to the K phase under deformation conditions. This transformation has been reported in previous investigations [9] to occur under deformation conditions in alloys of 4.6% Si subjected to rolling or tensile loading conditions. In the present investigation, cylindrical specimens of both silicon contents were subjected to dynamic loading in both the solution-treated and the aged conditions. In all cases the 4.37% Si alloy exhibited an expansion which was 2 - 3% greater than that of the 4.67% Si alloy. Under the loading conditions used in this study none of the alloys exhibited complete spallation, although incipient spall [2] and extensive microstructural damage were observed, as shown in Fig. 1 for a solution-treated 4.67% Si alloy specimen. This ability to limit the extent of internal fracture by spallation provided an opportunity to investigate the microstructural origins of spall fracture, and especially to examine the relationship between void formation and microstructural details. Some of the specimens exhibited complete separation over a limited region, thereby permitting the use of SEM techniques to examine the fracture surfaces.
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Fig. 1. Optical micrograph of a dynamically loaded Cu-4.67%Si alloy solution treated at 700 °C for 48 h.
3.1. Deformation structures resulting from dynamic loading In an earlier investigation [8] it was shown that the deformation microstructures generated by dynamic loading in copper-base alloys containing up to 3.3% Si are generally similar to those observed after quasi-static loading conditions [ 10]. Dislocation cells and tangles were observed in alloys containing less than 1% Si and bands of heavy dislocation density combined with twins were observed in alloys containing 1 - 2% Si. Alloys with higher silicon contents ranging up to 3.3% exhibited both twins and stacking faults. The 4.37% Si alloy of the present study revealed microstructural damage in the form of stacking fault configurations of even higher density than that observed in the lower silicon alloys; this behavior is consistent with the tendency for silicon to reduce the stacking fault energy as the solute content is increased. The Cu-4.67%Si alloy in the solutiontreated condition prior to loading exhibited stacking fault arrays, as shown in Fig. 2. X-ray diffraction analysis of the solution-treated and dynamically loaded Cu-4.67%Si alloy revealed diffraction peaks at values of 20 corresponding to 40.63 ° , 46.15 ° and 60.74 ° . These 20 values are in agreement with the expected peak positions for the strain-induced K phase [6]. The microstructure of this alloy after dynamic loading exhibited significant amounts of strain-induced ~ phase plates, as shown in Fig. 3. Previous work [9, 11] has shown that a similar strain-induced K phase is induced by quasi-static loading conditions in
Fig. 2. Bright field transmission electron micrograph showing the stacking faults in a solution-treated Cu-4.67%Si alloy.
alloys exceeding approximately 4.5% Si. A comparison of the typical micrographs in Figs. 2 and 3 lends support to earlier arguments [9] that the stacking fault configurations are probable nucleating sites for the formation of the strain-induced ~ phase under deformation conditions. Numerous defects resembling microvoids were also observed consistently in the thin foil microstructures of this alloy after dynamic loading; such defects are shown in Fig. 3 and are seen to be located at irregularities or steps along the interfaces separating the strain-induced K phase and the a matrix. Specimens of both silicon contents were subjected to aging treatments after solution treatment in order to develop a mixture of the a and K phases prior to dynamic loading. The ~ phase is in the form of platelets with a variable thickness, as is shown in Fig. 4 for a 4.67% Si alloy aged at 470 °C for 22 h. After
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t Fig. 3. TEM micrograph exhibiting strain-induced K phase plates produced by the dynamic loading of a Cu-4.67%Si alloy in the solution-treated condition. The void formation along 0~-K interfaces should be
dynamic loading, the microstructure exhibited a high density of deformation twins and arrays of fine spots associated preferentially with the interfaces, as shown in Fig. 5; these spots, like those in Fig. 3, are indicative of internal damage resulting from high strain rate loading and are interpreted as evidence of microvoids. Figures 6(a) and 6(b) provide a further indication of the nature of the damage resulting from dynamic loading. Figure 6(a) shows the severe distortion experienced by the K phase plates and the arrangement of voids along the interfacial regions. In Fig. 6(b) where the arrangement and interconnection of voids can be seen in greater detail there is evidence of interfacial separation due to the interconnection of closely spaced microvoids. Similar observations were noted in the microstructures of the Cu-4.37%Si alloy that was aged at 350 °C for 36 h to precipitate the K phase; as an example, Fig. 7 shows evidence of void development along the interfaces separating the K phase precipitates in the matrix. In all cases the microstructures exhibited some general distribution of voids in
noted.
Fig. 4. TEM micrograph showing the variable range of thickness of K plates formed during aging at 470 °C for 22 h in a Cu-4.67%Si alloy. The foil normal is close to the (110) direction.
F~. 5. TEM micrograph showing deformation twins in a dynamically loaded Cu-4.67%Si alloy aged at 470 °C for 22 h. The associated selected area electron diffraction pattern exhibits streaking. The foil orientation is close to the (110) direction.
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Fig. 7. An electron micrograph showing voids formed during the dynamic loading of a Cu-4.37%Si alloy aged at 350 °C for 36 h. (a)
(b) Fig. 6. (a) TEM micrograph showing typical microstructural damage, including void formation, after the shock loading of a Cu-4.67%Si alloy aged at 470 °C for 22 h; (b) TEM micrograph exhibiting void interconnection along ~-K interfaces due to the dynamic loading of a Cu-4.67%Si alloy aged at 470 °C for 22 h.
addition to the configurations associated with interfaces. The predominance of twinning in Cu-Si alloys subjected to dynamic loading is similar to the behavior observed in other f.c.c, metals and alloys [12 - 14]. The streaking of diffraction spots observed in the electron diffraction patterns of the present study is similar to that reported by Murr e t al. [15] in 304 stainless steel which also exhibited microtwinning after shock loading. In an earlier investigation of Cu-Si alloys [8] it was reported that the principal substructure effects resulting from dynamic loading were dislocation cells and tangles. The present work confirmed that, as the stacking fault energy is decreased with silicon content in excess of 3%, the deformation mechanism under dynamic loading conditions tends to be a twinning mode. The observation of a number of microstructures showed that the occurrence of twinning depends not only on a decrease in the stacking fault energy but also on the constraints imposed on the deformation process by the grain size of the parent phase; furthermore, the ~ plates that partition the a matrix provide an additional barrier to any deformation resulting from dislocation glide and/or
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stacking fault expansion. Other variables, including the peak pressure [16] and the pulse duration [ 17] associated with dynamic loading conditions, are expected to influence the occurrence of twinning. For example, when shock pressures exceeded 28 kbar, twins were observed in copper and Cu-A1 alloys [12] for which the stacking fault energy varied from 7.0 X 10 -2 to 0.25 × 10 -2 J m -2 depending on aluminum content. Such twins were observed in 70:30 brass [13] for which the stacking fault energy was estimated to be 1.4 X 10 -2 J m - 2 but only when the shock pressures exceeded 200 kbar. ' The microvoid features shown in Fig. 3 and Figs. 5 - 7 are similar to the defect structures produced during the radiation damage of materials. It has been suggested [18] that the shock hardening of stainless and ferritic steels may simulate radiation damage effects in these materials. Weertman [ 19] has predicted on the basis of dislocation theory that rapid loading could induce vacancy supersaturation; in addition, Kressel and Brown [20] have shown by means of electrical resistivity measurements that a high concentration of point defects, predominantly vacancies, is created in shock-loaded nickel. Vacancies or vacancy clusters have also been suggested as probable defects involved in the thermal recovery of shock-loaded and explosively formed Inconel 600 [ 2 1 ] , and it has been proposed that vacancies produced in high strain regions condense to form voids [22]. The formation of microvoids at the coherent boundary separating the a phase plates and the a matrix, as observed in the present work, is consistent with reports of voids at twin boundaries [ 23 - 25] and at the martensitematrix interface in Ti-6A1-4V alloys subjected to ductile fracture under static loading conditions [ 2 3 ] . The generation of linear arrays of voids as shown in Figs. 6(a) and 6(b) is also similar to the transmission electron microscopy (TEM) observations of microcracks at dislocation cell boundaries reported by Gardner et al. [ 2 5 ] . 3.2. Fractographic studies SEM techniques were employed to compare the fracture surface characteristics of Cu-Si alloys subjected to quasi-static and high strain rate loading conditions. The dynamically loaded cylinders exhibited many axial sur-
face cracks terminating in the radial direction below the potential spall plane. The existence of these cracks enabled SEM examination of the fracture surfaces produced by dynamic loading. The SEM observations were made with the viewing direction adjusted nearly normal to the macroscopic fracture surface. The fracture surfaces produced by dynamic loading of both the solution-treated and the aged specimens of the Cu-4.67%Si alloys exhibited characteristically ductile fracture modes, as revealed by the dimples and evidence of void interconnection shown in Figs. 8 and 9. In some cases the voids were associated with resolvable second-phase particles, b u t this was not a general observation. The fracture surfaces of the Cu-Si alloys after quasi-static loading exhibited a characteristic ductile failure mode in the form of elongated shear dimples as shown in Fig. 10 and the frequent occurrence of tear dimples. These features can be accounted for in terms of a fracture process in which microvoids elongate and coalesce under shear stress with the tear dimples being formed as a result of a tearing mode of microvoid coalescence; these
Fig. 8. SEM micrograph of the fracture surface p r o d u c e d by the dynamic loading of a solutiontreated C u - 4 . 6 7 % S i alloy. The void associated with an inclusion particle is s h o w n at A.
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Fig. 9. SEM photomicrograph showing dimple structure in a dynamically loaded Cu-4.67%Si alloy aged at 470 °C for 22 h.
Fig. 10. SEM fractograph of the lip region of the fracture surface of a Cu-4.67%Si alloy aged at 470 °C for 22 h and then pulled in tension.
features are in agreement with observations reported in precipitation-hardened alloys [26]. When the quasi-static fracture surfaces are compared with those resulting from dynamic loading, there is an order-of-magnitude difference in the average dimple size exhibited by the fracture surfaces as illustrated by Figs. 9 and 10. The fracture of plus K alloys under dynamic conditions results in an average dimple size of 0.4 pm, whereas the fracture of the same alloy with the same microstructure under quasi-static conditions results in an average dimple size of 7 urn. The relatively fine dimple sizes resulting from fracture under dynamic conditions, together with the low level of macroscopic strain prior to dynamic fracture, indicate that although fracture occurs by a ductile mode the failure process is controlled by the profuse nucleation of voids that can interconnect before reaching sizes that are resolvable by optical microscopy; this mechanism is consistent with the TEM observations of void distributions along -~ interfaces as shown in Figs. 5 and 6. The SEM fractographs from dynamically loaded solution-treated alloys revealed an average dimple size of 5.4 ~m, a value which is comparable with that observed in the statically fractured aged specimens. A comparison of Figs. 9 and 10 indicates that dynamic loading leads to more severe shear or tearing along the ridges separating the dimples along the fracture surfaces. The definite differences in the dimple sizes between the solution-treated and the aged specimens, as shown in Figs. 8 and 9, lead to the conclusion that the presence of the K phase promotes the nucleation of voids and enhances their interconnection, leading to fracture. By combining the results of TEM and SEM techniques it is possible to formulate a consistent description of the dynamic fracture process in two-phase Cu-Si alloys. The low energy coherent interfaces separating ~ plates from the ~ matrix provide an unexpectedly high density of sites for the nucleation of voids with an approximate spacing of 0.3 0.5 pm, as shown in Figs. 5 and 6. The close proximity of voids in planar configurations along the a-K interface provides the basis for initiating microcracks which can lead to specimen failure. Previous workers have contended that failure under dynamic conditions in polycrystalline metals is initiated by
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impurities, whereas in single crystals subgrain boundaries can serve as the initiation sites under static loading [ 25, 27]. It has also been shown in two-phase Ti-6A1-4V alloys that the process of ductile fracture is initiated by the p~eferential nucleation of voids at the interface between martensite plates and the a matrix; void growth in the titanium alloy appeared to depend on the plate lengths of the martensite phase [ 2 3 ] . Theresults of previous investigations indicate that the a-K interfaces of Cu-Si alloys are characterized by a high degree of morphological perfection as a result of an extremely low misfit 5 which was estimated to be 5.5 × 10 -4 [9, 28]. As a result of this.low degree of misfit, interfacial dislocations are expected to appear only rarely along the broad faces of the K plates. Kinsman et al. [28] have estimated the ledge spacing along a-K interfaces to be approximately 10 -4 cm on the basis of measurements of the rate of thickening of K plates at temperatures over the range 555 - 695 °C. When these ledge spacing values are compared with the average void spacings of approximately 0.3 - 0.5 pm along the ~-K interface, it is tempting to conclude that incoherent ledges along otherwise coherent interfaces can offer preferred sites for the nucleation of voids under dynamic loading conditions. The evidence of void formation at the coherent interfaces, as shown in Figs. 3, 5 and 6, adds positive support for such a mechanism. The dynamic fracture process in the two-phase Cu-Si alloys is controlled by the rate of microvoid nucleation along these interfaces; this conclusion follows from the dimple size reduction and the increase in microvoid density due to dynamic loading. ACKNOWLEDGMENTS
This work was supported in part by the National Science Foundation under Grant DMR-75-21843. Discussions with Professor I. M. Fyfe are gratefully acknowledged. REFERENCES 1 D. A. Shockey, L. Seaman and D. R. Curran, in R. W. Rohde, B. M. Butcher, J. R. Holland and C. H. Karnes (eds.), Metallurgical Effects at High Strain Rates, Plenum, New York, 1973, p. 473.
2 W. B. Jones and H. I. Dawson, in R. W. Rohde, B. M. Butcher, J. R. Holland and C. H. Karnes
3 4 5 6 7 8 9 10
11 12 13 14 15 16 17
18 19
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25 26 27 28
(eds.), Metallurgical Effects at High Strain Rates, Plenum, New York, 1973, p. 443. D. R. Curran, L. Seaman and D. A. Shockey, Phys. Today (January 1977) 46. S. E. Axter, W. B. Jones and D. H. Polonis, Metallography, 8 (1975) 425. A. Howie and P. R. Swann, Philos. Mag., 6 (1961) 1215. G. A. Dreyer and D. H. Polonis, Trans. AIME, "221 (1961) 1074. I. M. Fyfe and R. R. Ensminger, J. Mech. Phys. Solids, 14 (1966) 231. H. Vora, I. M. Fyfe and D. H. Polonis, Mater. Sci. Eng., 32 (1978) 129. M. B. Kasen, R. Taggart and D. H. Poionis, Trans. Am. Soc. Met., 60 (2) (1967) 144. P. R. Swann, in G. Thomas and J. Washburn (eds.), Electron Microscopy and the Strength of Crystals, Interscience, New York, 1963, p. 131. S. D. Dahlgren, W. F. Flanagan and D. H. Polonis, Trans. Metall. Soc. AIME, 2 (1966) 1071. O. Johari and G. Thomas, Acta Metall., 12 (1964) 1953. L. E. Murr and F. I. Grace, Exp. Mech., 5 (1969) 145. R. L. Nolder and G. Thomas, Acta Metall., 12 (1964) 227. L. E. Murr, J. V. Foltz and F. D. Aitman, Philos. Mag., 23 (1971) 1011. F. I. Grace, M. C. Inman and L. E. Murr, J. Phys. D, 1 (1968) 1437. L. E. Murr and K. P. Staudhammer, in J. V. Sanders and D. J. Goodchild (eds.), Electron Microscopy 1974, Proc. 8th Int. Congr. for Electron Microscopy, Vol. 1, Australian Institute of Metals, Canberra, 1974, p. 592. D. R. Ireland, M. Kargilaski and R. A. Wullaert, Trans. Am. Nucl. Soc., 13 (1970) 140. J. Weertman, in P. G. Shewmon and V. F. Zackay (eds.), Response of Metals to High Velocity Deformation, Interscience, New York, 1961, p. 205. H. Kressel and N. Brown, J. Appl. Phys., 38 (1967) 1618. L. E. Murr and B. H. Maddala, Mater. Sci. Eng., 7 (1971) 286: R. W. Bauer and H. G. F. Wilsdorf, Scr. Metall., 7 (1973) 1213. H. Margolin and Y. Mahajan, Metall. Trans. A, 9 (1978) 781. E. Roberts and P. G. Partridge, Deformation Twinning, Gordon and Breach, New York, 1964, p. 378. R. N. Gardner, T. C. Pollock and H. G. F. Wilsdorf, Mater. Sci. Eng., 29 (1977) 169. H. C. Rogers, Trans. AIME, 218 (1960) 498. J. J. Gilman and F. R. Tuler, Int. J. Fract. Mech., 6 (1970) 169. K. R. Kinsman, H. I. Aaronson and E. Eichen, Metall. Trans., 2 (1971) 1041.
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Deformation Substructures and Fracture B e h a v i o r o f C u - S i A l l o y s S u b j e c t e d to Dynamic Loading C o n d i t i o n s K. V. R A O and D. H. P O L O N I S
Metallurgical Engineering Division, University of Washington, Seattle, Wash. 98195 (U.S,A.) (Received D e c e m b e r 3, 1 9 7 9 )
SUMMARY
The microstructures and fracture surfaces of dynamically loaded Cu-Si alloys were examined by means of transmission electron microscopy and scanning electron microscopy techniques. The features of the deformation and fracture processes are interpreted in terms of the microstructural damage resulting from dynamic loading. The occurrence of a straininduced phase transformation in dynamically loaded solution-treated alloys was confirmed by the results o f both transmission electron microscopy and X-ray diffraction analysis. The deformation substructure after dynamic loading exhibits a profusion of microtwins in the matrix and the preferred nucleation of microvoids along ~-~ interfaces. A ductile fracture mode characterizes the failure under both static and dynamic loading conditions, but a lesser ductile mode associated with smaller dimple sizes was evident in dynamically loaded alloys that had been aged to produce the ~ phase.
1. I N T R O D U C T I O N
The response of a given material to dynamic loading conditions depends on a combination of factors including the stress history, the microstructure, the dislocation mobility and the presence of discontinuities and flaws. The prediction of dynamic failure in materials subjected to high intensity stress waves is limited by an inability to understand h o w the fracture process is controlled by changes in the microstructure and constitution of a material. Uncertainty still exists regarding the particular fracture mechanism
or combination of mechanisms that apply when the microstructure of a metal is modified by heat treatment or when the stress field conditions are changed. The complexity of the problem is illustrated by the processes of void formation and coalescence which are often important steps in the ductile fracture mechanism under both static and dynamic loading conditions. Previous workers have shown that impurity second-phase particles usually control the nucleaticn of voids leading to spallation under dynamic loading conditions [1 - 4]. In order to resolve the influence of precipitation processes on the dynamic behavior of materials, it is necessary to minimize the influence of impurity second-phase particles by employing alloys of high purity and carefully controlled composition. In the present work the response of Cu-Si alloys to dynamic loading conditions was investigated. The addition of silicon reduces the stacking fault energy of copper and promotes the formation of planar dislocation arrays and expanded faults [ 5]. Compositions with more than 4.6% Si (all alloy compositions are expressed in weight percentages) exhibit the formation of the h.c.p. K phase after appropriate thermal or mechanical processing [6]. The ~ phase forms as platelets parallel to {lll}~.c.c.: these platelets act as barriers for cross-slip processes and are expected to influence the flow and fracture behavior accompanying dynamic loading. In this paper we report on the microstructural changes and fracture modes due to the dynamic loading of Cu-Si alloys after solution treatment and after aging to produce a twophase mixture of the a and K phases. The influence of ~ platelets on the microstructural origins of the voids leading to ductile fracture was of special interest in this work.
16 2. EXPERIMENTAL PROCEDURE
Cu-Si alloys containing 4.37% Si and 4.67% Si were prepared by vacuum melting and casting from OFHC copper and 99.999% Si. The cylindrical rods were homogenized at 750 °C for 48 h and subsequently cold swaged to rods 1.3 cm in diameter with intermediate annealing treatments. Hollow cylindrical test specimens with an inside diameter of 5 mm, an outside diameter of 10 mm and a length of 19.1 mm were machined from the homogenized rods and solution treated at 700 °C for 48 h. Aging treatments were conducted at 470 and 350 °C on the solution-treated and quenched specimens. Prior to heat treatment the specimens were encapsulated in Vycor tubing, evacuated and backfilled with a partial pressure of argon. An exploding wire system developed by Fyfe and Ensminger [7] was employed to generate the dynamic loading conditions in the hollow cylindrical specimens. The system consists of two 15 p F capacitors charged to a nominal 20 kV and providing an energy storage of 3000 J per capacitor for a total of 6000 J. Specimens loaded in this fashion expand in the radial direction with complete symmetry and with little variation along the axis of the cylinder. This loading technique results in a biaxial strain condition in thickwalled cylinders where the hoop and radial strains are non-zero. The high intensity stress waves required for the dynamic loading of thick-walled cylinders are generated by discharging the capacitors through a polyethylene-coated wire. The polyethylene is driven against the inside diameter of the specimen and acts as a cylindrical piston. The pressure pulse has a rise time of approximately 0.1 ps, and the maximum load can be as high as 25 kbar. The present experiments employed copper wires 0.84 mm in diameter encased in a polyethylene jacket 2.95 mm in diameter. The static tensile tests employed an Instron testing machine operated at a strain rate of 3.3 × 10 -4 s-1 at room temperature. The X-ray diffraction analyses were performed with a Philips vertical diffractometer and Cu Ks radiation generated at 40 kW and 20 mA. A graphite single-crystal m o n o c h r o m a t o r was used to discriminate the Cu Ks radiation. Thin foil specimens for transmission electron microscopy were prepared from slices
10 × 10 -3 in thick cut from the inner surface of the cylinders in a direction tangential to the circumference. The slices were chemically polished in a solution of 50% nitric acid, 25% acetic acid and 25% orthophosphoric acid. Discs 3.05 mm in diameter punched from these slices were electropolished in a jet polisher using a solution containing one part of nitric acid and two parts of methyl alcohol at temperatures below --20 °C. The thin foils were examined in a JEM-7 electron microscope operated at 100 kV. The examination of fracture surfaces was conducted by scanning electron microscopy (SEM) techniques.
3. RESULTS AND DISCUSSION
The 4.37% Si and 4.67% Si alloys used in the present study exhibited less macroscopic strain and a reduced Susceptibility to spallation compared with alloys of lower silicon content examined in an earlier study [8]. The Cu-4.37%Si alloy is close to the limit above which the supersaturated ~ phase in solutiontreated alloys is partially transformed to the K phase under deformation conditions. This transformation has been reported in previous investigations [9] to occur under deformation conditions in alloys of 4.6% Si subjected to rolling or tensile loading conditions. In the present investigation, cylindrical specimens of both silicon contents were subjected to dynamic loading in both the solution-treated and the aged conditions. In all cases the 4.37% Si alloy exhibited an expansion which was 2 - 3% greater than that of the 4.67% Si alloy. Under the loading conditions used in this study none of the alloys exhibited complete spallation, although incipient spall [2] and extensive microstructural damage were observed, as shown in Fig. 1 for a solution-treated 4.67% Si alloy specimen. This ability to limit the extent of internal fracture by spallation provided an opportunity to investigate the microstructural origins of spall fracture, and especially to examine the relationship between void formation and microstructural details. Some of the specimens exhibited complete separation over a limited region, thereby permitting the use of SEM techniques to examine the fracture surfaces.
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Fig. 1. Optical micrograph of a dynamically loaded Cu-4.67%Si alloy solution treated at 700 °C for 48 h.
3.1. Deformation structures resulting from dynamic loading In an earlier investigation [8] it was shown that the deformation microstructures generated by dynamic loading in copper-base alloys containing up to 3.3% Si are generally similar to those observed after quasi-static loading conditions [ 10]. Dislocation cells and tangles were observed in alloys containing less than 1% Si and bands of heavy dislocation density combined with twins were observed in alloys containing 1 - 2% Si. Alloys with higher silicon contents ranging up to 3.3% exhibited both twins and stacking faults. The 4.37% Si alloy of the present study revealed microstructural damage in the form of stacking fault configurations of even higher density than that observed in the lower silicon alloys; this behavior is consistent with the tendency for silicon to reduce the stacking fault energy as the solute content is increased. The Cu-4.67%Si alloy in the solutiontreated condition prior to loading exhibited stacking fault arrays, as shown in Fig. 2. X-ray diffraction analysis of the solution-treated and dynamically loaded Cu-4.67%Si alloy revealed diffraction peaks at values of 20 corresponding to 40.63 ° , 46.15 ° and 60.74 ° . These 20 values are in agreement with the expected peak positions for the strain-induced K phase [6]. The microstructure of this alloy after dynamic loading exhibited significant amounts of strain-induced ~ phase plates, as shown in Fig. 3. Previous work [9, 11] has shown that a similar strain-induced K phase is induced by quasi-static loading conditions in
Fig. 2. Bright field transmission electron micrograph showing the stacking faults in a solution-treated Cu-4.67%Si alloy.
alloys exceeding approximately 4.5% Si. A comparison of the typical micrographs in Figs. 2 and 3 lends support to earlier arguments [9] that the stacking fault configurations are probable nucleating sites for the formation of the strain-induced ~ phase under deformation conditions. Numerous defects resembling microvoids were also observed consistently in the thin foil microstructures of this alloy after dynamic loading; such defects are shown in Fig. 3 and are seen to be located at irregularities or steps along the interfaces separating the strain-induced K phase and the a matrix. Specimens of both silicon contents were subjected to aging treatments after solution treatment in order to develop a mixture of the a and K phases prior to dynamic loading. The ~ phase is in the form of platelets with a variable thickness, as is shown in Fig. 4 for a 4.67% Si alloy aged at 470 °C for 22 h. After
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t Fig. 3. TEM micrograph exhibiting strain-induced K phase plates produced by the dynamic loading of a Cu-4.67%Si alloy in the solution-treated condition. The void formation along 0~-K interfaces should be
dynamic loading, the microstructure exhibited a high density of deformation twins and arrays of fine spots associated preferentially with the interfaces, as shown in Fig. 5; these spots, like those in Fig. 3, are indicative of internal damage resulting from high strain rate loading and are interpreted as evidence of microvoids. Figures 6(a) and 6(b) provide a further indication of the nature of the damage resulting from dynamic loading. Figure 6(a) shows the severe distortion experienced by the K phase plates and the arrangement of voids along the interfacial regions. In Fig. 6(b) where the arrangement and interconnection of voids can be seen in greater detail there is evidence of interfacial separation due to the interconnection of closely spaced microvoids. Similar observations were noted in the microstructures of the Cu-4.37%Si alloy that was aged at 350 °C for 36 h to precipitate the K phase; as an example, Fig. 7 shows evidence of void development along the interfaces separating the K phase precipitates in the matrix. In all cases the microstructures exhibited some general distribution of voids in
noted.
Fig. 4. TEM micrograph showing the variable range of thickness of K plates formed during aging at 470 °C for 22 h in a Cu-4.67%Si alloy. The foil normal is close to the (110) direction.
F~. 5. TEM micrograph showing deformation twins in a dynamically loaded Cu-4.67%Si alloy aged at 470 °C for 22 h. The associated selected area electron diffraction pattern exhibits streaking. The foil orientation is close to the (110) direction.
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Fig. 7. An electron micrograph showing voids formed during the dynamic loading of a Cu-4.37%Si alloy aged at 350 °C for 36 h. (a)
(b) Fig. 6. (a) TEM micrograph showing typical microstructural damage, including void formation, after the shock loading of a Cu-4.67%Si alloy aged at 470 °C for 22 h; (b) TEM micrograph exhibiting void interconnection along ~-K interfaces due to the dynamic loading of a Cu-4.67%Si alloy aged at 470 °C for 22 h.
addition to the configurations associated with interfaces. The predominance of twinning in Cu-Si alloys subjected to dynamic loading is similar to the behavior observed in other f.c.c, metals and alloys [12 - 14]. The streaking of diffraction spots observed in the electron diffraction patterns of the present study is similar to that reported by Murr e t al. [15] in 304 stainless steel which also exhibited microtwinning after shock loading. In an earlier investigation of Cu-Si alloys [8] it was reported that the principal substructure effects resulting from dynamic loading were dislocation cells and tangles. The present work confirmed that, as the stacking fault energy is decreased with silicon content in excess of 3%, the deformation mechanism under dynamic loading conditions tends to be a twinning mode. The observation of a number of microstructures showed that the occurrence of twinning depends not only on a decrease in the stacking fault energy but also on the constraints imposed on the deformation process by the grain size of the parent phase; furthermore, the ~ plates that partition the a matrix provide an additional barrier to any deformation resulting from dislocation glide and/or
20
stacking fault expansion. Other variables, including the peak pressure [16] and the pulse duration [ 17] associated with dynamic loading conditions, are expected to influence the occurrence of twinning. For example, when shock pressures exceeded 28 kbar, twins were observed in copper and Cu-A1 alloys [12] for which the stacking fault energy varied from 7.0 X 10 -2 to 0.25 × 10 -2 J m -2 depending on aluminum content. Such twins were observed in 70:30 brass [13] for which the stacking fault energy was estimated to be 1.4 X 10 -2 J m - 2 but only when the shock pressures exceeded 200 kbar. ' The microvoid features shown in Fig. 3 and Figs. 5 - 7 are similar to the defect structures produced during the radiation damage of materials. It has been suggested [18] that the shock hardening of stainless and ferritic steels may simulate radiation damage effects in these materials. Weertman [ 19] has predicted on the basis of dislocation theory that rapid loading could induce vacancy supersaturation; in addition, Kressel and Brown [20] have shown by means of electrical resistivity measurements that a high concentration of point defects, predominantly vacancies, is created in shock-loaded nickel. Vacancies or vacancy clusters have also been suggested as probable defects involved in the thermal recovery of shock-loaded and explosively formed Inconel 600 [ 2 1 ] , and it has been proposed that vacancies produced in high strain regions condense to form voids [22]. The formation of microvoids at the coherent boundary separating the a phase plates and the a matrix, as observed in the present work, is consistent with reports of voids at twin boundaries [ 23 - 25] and at the martensitematrix interface in Ti-6A1-4V alloys subjected to ductile fracture under static loading conditions [ 2 3 ] . The generation of linear arrays of voids as shown in Figs. 6(a) and 6(b) is also similar to the transmission electron microscopy (TEM) observations of microcracks at dislocation cell boundaries reported by Gardner et al. [ 2 5 ] . 3.2. Fractographic studies SEM techniques were employed to compare the fracture surface characteristics of Cu-Si alloys subjected to quasi-static and high strain rate loading conditions. The dynamically loaded cylinders exhibited many axial sur-
face cracks terminating in the radial direction below the potential spall plane. The existence of these cracks enabled SEM examination of the fracture surfaces produced by dynamic loading. The SEM observations were made with the viewing direction adjusted nearly normal to the macroscopic fracture surface. The fracture surfaces produced by dynamic loading of both the solution-treated and the aged specimens of the Cu-4.67%Si alloys exhibited characteristically ductile fracture modes, as revealed by the dimples and evidence of void interconnection shown in Figs. 8 and 9. In some cases the voids were associated with resolvable second-phase particles, b u t this was not a general observation. The fracture surfaces of the Cu-Si alloys after quasi-static loading exhibited a characteristic ductile failure mode in the form of elongated shear dimples as shown in Fig. 10 and the frequent occurrence of tear dimples. These features can be accounted for in terms of a fracture process in which microvoids elongate and coalesce under shear stress with the tear dimples being formed as a result of a tearing mode of microvoid coalescence; these
Fig. 8. SEM micrograph of the fracture surface p r o d u c e d by the dynamic loading of a solutiontreated C u - 4 . 6 7 % S i alloy. The void associated with an inclusion particle is s h o w n at A.
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Fig. 9. SEM photomicrograph showing dimple structure in a dynamically loaded Cu-4.67%Si alloy aged at 470 °C for 22 h.
Fig. 10. SEM fractograph of the lip region of the fracture surface of a Cu-4.67%Si alloy aged at 470 °C for 22 h and then pulled in tension.
features are in agreement with observations reported in precipitation-hardened alloys [26]. When the quasi-static fracture surfaces are compared with those resulting from dynamic loading, there is an order-of-magnitude difference in the average dimple size exhibited by the fracture surfaces as illustrated by Figs. 9 and 10. The fracture of plus K alloys under dynamic conditions results in an average dimple size of 0.4 pm, whereas the fracture of the same alloy with the same microstructure under quasi-static conditions results in an average dimple size of 7 urn. The relatively fine dimple sizes resulting from fracture under dynamic conditions, together with the low level of macroscopic strain prior to dynamic fracture, indicate that although fracture occurs by a ductile mode the failure process is controlled by the profuse nucleation of voids that can interconnect before reaching sizes that are resolvable by optical microscopy; this mechanism is consistent with the TEM observations of void distributions along -~ interfaces as shown in Figs. 5 and 6. The SEM fractographs from dynamically loaded solution-treated alloys revealed an average dimple size of 5.4 ~m, a value which is comparable with that observed in the statically fractured aged specimens. A comparison of Figs. 9 and 10 indicates that dynamic loading leads to more severe shear or tearing along the ridges separating the dimples along the fracture surfaces. The definite differences in the dimple sizes between the solution-treated and the aged specimens, as shown in Figs. 8 and 9, lead to the conclusion that the presence of the K phase promotes the nucleation of voids and enhances their interconnection, leading to fracture. By combining the results of TEM and SEM techniques it is possible to formulate a consistent description of the dynamic fracture process in two-phase Cu-Si alloys. The low energy coherent interfaces separating ~ plates from the ~ matrix provide an unexpectedly high density of sites for the nucleation of voids with an approximate spacing of 0.3 0.5 pm, as shown in Figs. 5 and 6. The close proximity of voids in planar configurations along the a-K interface provides the basis for initiating microcracks which can lead to specimen failure. Previous workers have contended that failure under dynamic conditions in polycrystalline metals is initiated by
22
impurities, whereas in single crystals subgrain boundaries can serve as the initiation sites under static loading [ 25, 27]. It has also been shown in two-phase Ti-6A1-4V alloys that the process of ductile fracture is initiated by the p~eferential nucleation of voids at the interface between martensite plates and the a matrix; void growth in the titanium alloy appeared to depend on the plate lengths of the martensite phase [ 2 3 ] . Theresults of previous investigations indicate that the a-K interfaces of Cu-Si alloys are characterized by a high degree of morphological perfection as a result of an extremely low misfit 5 which was estimated to be 5.5 × 10 -4 [9, 28]. As a result of this.low degree of misfit, interfacial dislocations are expected to appear only rarely along the broad faces of the K plates. Kinsman et al. [28] have estimated the ledge spacing along a-K interfaces to be approximately 10 -4 cm on the basis of measurements of the rate of thickening of K plates at temperatures over the range 555 - 695 °C. When these ledge spacing values are compared with the average void spacings of approximately 0.3 - 0.5 pm along the ~-K interface, it is tempting to conclude that incoherent ledges along otherwise coherent interfaces can offer preferred sites for the nucleation of voids under dynamic loading conditions. The evidence of void formation at the coherent interfaces, as shown in Figs. 3, 5 and 6, adds positive support for such a mechanism. The dynamic fracture process in the two-phase Cu-Si alloys is controlled by the rate of microvoid nucleation along these interfaces; this conclusion follows from the dimple size reduction and the increase in microvoid density due to dynamic loading. ACKNOWLEDGMENTS
This work was supported in part by the National Science Foundation under Grant DMR-75-21843. Discussions with Professor I. M. Fyfe are gratefully acknowledged. REFERENCES 1 D. A. Shockey, L. Seaman and D. R. Curran, in R. W. Rohde, B. M. Butcher, J. R. Holland and C. H. Karnes (eds.), Metallurgical Effects at High Strain Rates, Plenum, New York, 1973, p. 473.
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