Some SEM observations on high-temperature fracture in aluminum

METALLOGRAPHY 10, 135-148 (1977)

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Some SEM Observations on High-Temperature Fracture in Aluminum V. KUTUMBA RAO, VAKIL SINGH, and P. RAMA RAO Department of Metallurgical Engineering, Banaras Hindu University, Varanasi-221005, India

A scanning electron microscope has been employed to study fracture in AI under creep and low cycle fatigue conditions at high temperatures. For purposes of comparison, fracture of Al at room temperature in tension and low cycle fatigue has been examined. The conclusion emerges that fracture mode of A1 during creep at 548 K (0.59 Tin, where Tm is the melting temperature in degrees Kelvin) is analogous to tensile fracture at the ambient temperature. On the other hand, under low cycle fatigue loading, fracture at the higher temperature, namely 648 K (0.69 Tin), is significantly different with fatigue, ductile, and intergranular rupture components constituting a mixed fracture mode in comparison to the almost-complete fatigue fracture at room temperature. Some details concerning the creep and high-temperature fatigue fracture processes have also been noted.

Introduction The advent of the scanning electron microscope (SEM) has placed in the hands of fractographers in recent years the most useful tool yet for the study of fracture surfaces. Although considerable use of this tool has been made in a general way for more than a decade, its purposeful utilization for studying high-temperature fracture is comparatively recent [1-3]. The present work is concerned with the use of S E M to delineate features of fracture of A1 under creep as well as low cycle fatigue conditions at high temperatures. Observations of fracture at r o o m temperature under similar loading conditions have been made for purposes of comparison.

Experimental E C grade AI o f 99.76% purity (containing by wt% 0.015 Fe and 0.10 Si as impurities) was used in room-temperature tensile and high © Elsevier North-Holland, Inc., 1977

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temperature creep tests. This grade of A1 as also superpurity A1 of 99.99% purity (containing by wt% 0.0001 Fe, 0.0015 Si, and 0.005 Cu as impurities) were used in low cycle fatigue tests. Round specimens (3.6 mm diameter and 25 mm gauge length) with threaded ends were used in constant-load creep tests that were carried out at 548 K (0.59 Tr,) at two initial stress levels of 14.8 and 17.7 MN m -2 on an Amsler Type STFM 746 Creep machine. Standard Hounsfied No. 12 specimens were used for tensile testing on a modified Hounsfield Tensometer [4]. Push-pull low cycle fatigue tests were conducted at 648 K (0.69 Tin) as also at room temperature (302 K, 0.32 Tin) on waisted specimens (14.5 mm diameter and 4.8 mm gauge length) using an Instron Universal Testing Machine suitably modified for this purpose. Extension cycle control with varying total strain amplitude in the range --- 1 to ___ 10% and a cycle frequency of 5 cpm were used in the low cycle fatigue tests. In all cases, fracture surfaces were examined in a Cambridge Stereoscan S-2 scanning electron microscope using the secondary emission mode of operation.

Results and Discussion (~ CREEP F R A C T U R E In Table 1 the test conditions employed in creep are summarized along with the resulting relevant creep data. Figure 1 shows SEM pictures of specimens fractured in creep at 548 K under a stress of 14.8 MN m -z and in simple tension at 302 K and a strain rate of 1.8 x 10-1 hr -1. In both cases the fracture surface shows a number of essentially

TABLE 1 Creep Properties of EC Grade A1 Tested at 548 K

Stress (MN. m -2)

Minimum Creep Rate (hr -1)

Elongation at Fracture (%)

Time to Rupture (hr)

14.8 17.7

1.89 x 10-a 5.35 × 10-a

17.5 32.0

117 42

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(a)

(b) FIG. 1. SEM micrographs of fracture surface of EC grade A1 (a) tested in creep at 548 K and 14.8 MN m -2 (b) tested in simple tension at room temperature (302 K).

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(a)

(b) FIG. 2 (a) and (b). SEM micrographs of fracture surface of EC grade A1 tested in creep at 548 K and 14.8 MN m -~.

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equiaxed microvoids 1, which are typical of ductile transgranular rupture. No evidence could be seen for intergranular fracture at 548 K (0.59 Tin) which lies in the homologous temperature range in which most metals and alloys fail by intergranular cracking. This is in accord with the observation by Grant and co-workers [5, 6] that during creep of A1, transgranular fracture persists at high temperatures owing at least partly to extensive grain-boundary migration which prevents or at least delays the formation and growth of cracks on grain boundaries. An interesting observation in the present study (Fig. 2) concerns the growth process of the voids of Fig. 1. Figure 2(a) shows that the area around each of these voids is composed of a number of smaller voids; we shall refer to the former as "primary" voids and to the latter as "secondary" voids. A careful examination of this area shows that the secondary voids coalesce by stretching out under the action of the tensile stress until the intervoid material fractures. Those secondary voids that are immediately on the periphery of the primary void would be evidently the first to merge in this manner, because the area of effective cross section here is the smallest. The eventual result of this process is that the primary void grows in size until it merges with a neighboring primary void which has also been growing independently in a similar manner [Fig. 2(b)]. Thus the process of linking-up of the microvoids in the "void-sheet" mechanism may involve not a simple plastic tearing of the intervoid space as is generally believed, but the coalescence of a number of smaller voids surrounding each bigger void. The effect of increasing the applied stress on the above process of fracture appears to be a reduction in the size of the primary voids (Fig. 3). Also, there are comparatively fewer secondary voids and their structure is less well resolved. This is presumably a result of the much more rapid coalescence of the secondary voids under the action of a higher stress, and thus a higher strain rate (Table 1), while the overall short duration of the test itself accounts for the smaller size of the general population of primary voids. The very small primary void size and the near absence of resolvable secondary voids in specimens fractured at room temperature [Fig. l(b)] at a strain rate nearly two orders of magnitude higher than the strain rates obtained during creep at 548 K is in accordance with the above argument. It is possible that the strain rate sensitivity (m) of flow stress has an influence on the fracture 1Usually called "dimples" owing to their appearance as such in replicas examined by transmission electron microscopy; however, the general and more accurate term "voids" (or "microvoids") is preferred by the present authors.

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(a)

(b) Fl6.3 (a) and (b). S E M micrographs of fracturesurface of E C grade A I tested in creep at 548 K and 17.7 M N m -2.

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TABLE 2 Fracture Life of AI Tested in Low Cycle Fatigue Fracture Life (Number of Cycles) Nominal Strain Amplitude (%)

EC Grade A1 at 302 K

648 K

Superpurity Al at 648 K

+- 1 -+2 -+4 -+ l0

-4,365 -37

22,300 8,350 2,480 640

-17,000 5,100 1,250

process described above [7], and it is relevant to note that m is lower at lower temperatures and higher strain rates. (b) L O W C Y C L E F A T I G U E F R A C T U R E In Table 2 the test conditions employed in low cycle fatigue experiments and the relevant results are summarized. The variation of fatigue life with the total strain range in both grades of A1 is in accordance with the well-known Coffin-Manson law [8] applicable to low cycle fatigue. The fatigue lives in superpurity A1 are significantly higher than in EC grade A1 at all strain amplitudes tested. This can be explained in terms of observed higher grain-boundary mobility in superpurity A1 as evident from larger initial load drops in the load response curves and more rapid grain-boundary configurational changes as seen in the surface microstructures [9]. SEM pictures of fracture surfaces (Fig. 4) show that the fatigue failure at 648 K (0.69 Tin) in the present study is a mixture of transgranular and intergranular modes. The micrographs of Fig. 4 show intergranular cracks, Zapffe and Worden striations characteristic of fatigue crack propagation, and microvoids associated with ductile rupture. The noticeably cleaner appearance of the fracture surface of superpurity A1 [Fig. 4 (a)] compared to EC grade A1 [Fig. 4(b)] may be attributed to a reduction in complicating features such as inclusions and second-phase particles in the higher-purity metal. The fact that some amount of

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(a)

(b) FIG. 4. SEM micrographs of fracture surface of specimens tested in low cycle fatigue at 648 K. (a) Superpurity A1, strain amplitude -+ 4%, (b) EC grade AI, strain amplitude -+ 1%.

Fracture

Surfaces in Aluminum

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(a)

(b) Fio. 5. S E M micrographs of fracture surface of specimens tested in low cycle fatigue at 648 K. (a) Superpurity Al, strainamplitude ± 2%, (b) E C grade AI, strainamplitude ± 10%.

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intergranular fracture is obtained at all during high-temperature fatigue of A1 is remarkable since, as has been shown in the present study as well as by others [5, 6], fracture of A1 during creep is always transgranular irrespective of temperature and strain rate. Blucher and Grant [10], using optical microscopy, have observed that fatigue fracture of A1 at temperatures greater than about 640 K is almost completely intergranular. However, in the present study, where the temperature of testing (648 K) was purposely chosen to verify whether intergranular fracture dominates during high-temperature fatigue, a careful examination of a number of micrographs reveals that a mixed fracture mode is obtained which is influenced by the strain amplitude. Figure 5, which shows fracture surfaces of specimens failed in fatigue at 648 K at strain amplitudes of -+2% [Fig. 5(a)] and -+10% [Fig. 5(b)], illustrates the observation that with increasing strain amplitude the fracture becomes increasingly transgranular, that is, the fracture increasingly resembles that obtained at room temperature [see Fig. 8(a)]. This is not surprising in view of the frequently observed fact [11, 12] that at high strain rates, even at high temperatures, the nature of deformation and fracture is quite similar to that at lower temperatures. In the present case, frequency of cycling being constant, strain rate increases with increasing strain amplitude. Observations of specimen surface clearly emphasize the fact that crack nucleation generally is intergranular (Fig. 6). Although slip bands are formed on the surface (Fig. 6), no slip band cracking is evident. Also seen in Fig. 6(a) are apparently "diamond"-shaped grains, a morphologic feature often encountered during fatigue deformation at high temperatures [10, 13]. The cracks formed at grain boundaries on the surface propagate, however, both trangranularly and intergranularly, as indicated by the occurrence of the familiar fatigue striations in grain interiors [Fig. 7(a)] as well as on grain facets [Fig. 7(b)]. In Fig. 8 micrographs of specimens fractured during fatigue at room temperature are recorded. The fracture here is completely transgranular, showing that the mixed mode of fracture at 648 K described above is a result of different fracture mechanisms operating at the higher temperature. The "stepped" appearance of the fracture surface [Fig. 8(a)] is a result of the propagation of cracks on planes at different levels [14]. Further, the fatigue striations here [Fig. 8(b)] are seen to be sharper and straighter than the ones observed during high-temperature fatigue (Fig. 7). The reason for this difference should lie in the difference in strength, plasticity, and the fracture life (Table 2) of the material at the two temperatures.

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(a)

(b) FIG. 6. SEM micrographs of surface of specimens tested in low cycle fatigue at 648 K and a strain amplitude of -+ 4%. (a) Superpurity AI, (b) EC grade A1.

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(a)

(b) FIG. 7. SEM micrographs of fracture surface of specimens tested in low cycle fatigue at 648 K. (a) Superpurity A1, strain amplitude -+ 2%, (b) EC grade A1, strain amplitude -+ 4%.

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FXG. 8 (a) and (b). SEM micrographs of fracture surface of specimens tested in low cycle fatigue at room temperature (302 K) at a strain amplitude of +- 2%.

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Summary 1. Scanning electron microscopy of fracture of A1 during hightemperature deformation has shown that fracture in creep at 548 K was typically ductile transgranular, while in low cycle fatigue at 648 K it was a mixture of transgranular and intergranular modes, the former increasing with increasing strain amplitude. 2. The ductile rupture process in creep involved the formation of "primary" microvoids and their growth by the consumption of a number of smaller peripheral "secondary" voids. 3. The high-temperature fatigue fracture process involved the nucleation of cracks generally at grain boundaries on the surface and their propagation through grain interiors as well as grain boundaries. We are grateful to Professor T. R. Anantharaman for provision o f facilities and encouragement. We thank Professor D. M. R. Taplin for his interest in this work and Mr. H. Kamler for assistance in the S E M work. One o f us (V. K. R.) was supported by a National Research Council (Canada)grant to spend a short tenure at the University o f Waterloo, Waterloo, Ontario, during which the SEM work was carried out.

References 1. A. E. Carden, A. J. McEvily, and C. H. Wells (Eds.), Fatigue at Elevated Temperatures, STP No. 520, ASTM (1973). 2. R. P. Simpson, G. J. Dooley, and T. W. Haas, Met. Trans. 5, 585 (1974). 3. B. J. Cane and G. W. Greenwood, Met. Sci. 9, 55 (1975). 4. V. Kutumba Rao, D. M. R. Taplin, and P. Rama Ran, Trans. Ind. Inst. Metals 23, 61 (1970). 5. I. S. Servi and N. J. Grant, J. Metals 3, 909 (1951). 6. A. W. Mullendore and N. J. Grant, Trans. A I M E 227, 319 (1963). 7. R. G. Fleck, C. J. Beevers, and D. M. R. Taplin, Metal Sci. 9, 49 (1975). 8. J. F. Taverneili and L. F. Coffin, Trans. A S M 51, 438 (1959). 9. Vakil Singh, P h . D . Thesis, Banaras Hindu University (1974). 10. J. T. Blucher and N. J. Grant, Trans. A I M E 239, 805 (1967). 11. J. C. Grosskreutz, Proc. Tenth Sagamore Army Materials Research Conference (J. J. Burke, N. L. Reed, and V. Weiss, Eds.), Syracuse University Press, Syracuse (1964). 12. S. Taira, High Temperature Structure and Materials (A. M. Freudenthal, B. A. Boley, and H. Liebowitz, Eds.), MacMillan, New York (1964). 13. Vakil Singh, P. Rama Rao, G. J. Cocks, and D. M. R. Taplin, J. Mat. Sci. (in press). 14. C. Laird and G. C. Smith, Phil. Mag. 7, 847 (1962). 15. A. J. Perry, J. Mat. Sci. 9, 1016 (1974). Received October, 1975