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Mechanical behavior and microstructural evolution upon annealing of the accumulative roll-bonding (ARB) processed Al alloy 1100
Materials Science and Engineering A 480 (2008) 148–159
Mechanical behavior and microstructural evolution upon annealing of the accumulative roll-bonding (ARB) processed Al alloy 1100 Charles Kwan a , Zhirui Wang a,∗ , Suk-Bong Kang b a
b
Department of Materials Science and Engineering, University of Toronto, 184 College Street, Toronto, Ont., Canada M2J 1L2 Materials Engineering Department, Korea Institute of Machinery and Materials, 66 Sangnam, Kyungnam, Changwon 641-010, Republic of Korea Received 17 April 2007; received in revised form 27 June 2007; accepted 3 July 2007
Abstract The results of mechanical testing on ultra-fine grained aluminum processed by accumulative roll-bonding (ARB) were analyzed with TEM observation and in accordance of the microstructural evolution upon annealing. It was found that rapid grain growth, with the corresponding decrease in strength, did not occur until the annealing temperature of 200 ◦ C or higher. The oxide rolled into the material near the bonding interfaces was seen to act as an obstruction for grain boundary migration across said interfaces. More interestingly, the strain near the interface due to the surface preparation technique used during ARB was found to form discontinuous segregates consisting of smaller grains formed during annealing or even ARB processing of higher number of cycles. Such phenomenon is attributed to recovery or polygonization due to the strain incurred. This study has also demonstrated that yield point phenomenon may be observed in a commercially pure fcc metal when the grain size is within a certain range. © 2007 Elsevier B.V. All rights reserved. Keywords: Accumulative roll-bonding; Aluminum; Mechanical behavior; Annealing; Microstructure evolution; Strain hardening
1. Introduction Materials with ultra-fine grain (UFG) structure (having a mean grain size between 100 nm and 1 m) have been gathering the interest of scientists for more than 50 years and especially in the past 20 years. Not only do the UFGed materials display increase in strength through Hall–Petch strengthening, but also these materials have high strain rate super plasticity [1–3] and low temperature super plasticity [4]. It has been shown that UFGed structure can be obtained through severe plastic deformation, methods such as torsion straining and equal channel angular pressing (ECAP). Most of these available processes lack the potential for a cost efficient scale up for the production of bulk UFGed material in large quantity. Saito et al. [5,6] have successfully created a method to form bulk UFGed structure through a processing technique called accumulative roll-bonding (ARB). The ARB process is essentially a process of multiple rolling, whereby after each roll to a 50% reduction in the sheet thickness, the sheet is cut into halves, surface treated and stacked up before
∗
Corresponding author. Tel.: +1 416 978 4412; fax: +1 416 978 4415. E-mail address: [email protected] (Z. Wang).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.07.022
it is rolled again [5,6]. By stacking the two halves between each 50% reduction roll, the thickness of the sheet remains unchanged throughout any given number of cycles. This process has been successfully applied to make UFGed sheets of aluminum and aluminum alloys [7–15], and IF steel [15,16]. Hall–Petch strengthening is based on the increase in the increased grain boundary volume fraction acting as obstacles for dislocation glide. Through the use of the ARB process, the tensile strength of 1100 aluminum reached 270 MPa after 6 cycles of ARB, corresponding to an increase of approximately 3 times from the 90 MPa tensile strength of the base material [14,15,17]. On the contrary, the elongation to failure decreases from approximately 40% to 5% upon even 1 cycle of ARB [14]. Although H¨oppel et al. have shown an increase of elongation to failure as the number of ARB cycles increases [17], the results presented by them strongly support the need of post-ARB annealing in order to recover at least some ductility. The microstructure resulting from the ARB process is shown to be in a pancake-like shaped grains surrounded by well-defined high angle boundaries [8–16]. In recent years, specifically, Huang et al. [13], Tsuji et al. [18], and Ueji et al. [19], showed that the lamellar boundaries in ARB materials consisted mostly of high angle boundaries, whereas the interconnecting
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boundaries between the pancake grains (running vertical through thickness) consisted of a mixture of low and high angle boundaries. Obviously, such high angle boundaries parallel to the rolling direction may bring on new and unique mechanical behavior to these materials. In addition, the ARB process also introduces layer interfaces within the bulk, and for materials that readily form oxide film, such as aluminum, the layer interfaces introduce particles (alumina in this case) into the bulk, which should also alter the mechanical behavior as well as the grain growth behavior. Up to date, there has been a lack of systematic studies on the mechanical behavior of ARB materials, and even less is known about the changes of microstructure as well as mechanical behavior upon annealing of this type of materials. Hence, the present study on the subject matter has been carried out.
three or more distinctive micrographs were measured to obtain the average grain size. Grain size of 1, 4, and 6 cycles samples upon annealing at 300 ◦ C or higher could not be obtained this way, as the grain size under these annealing conditions was too large to contain more than one or two grains in single TEM micrographs. Micrographs after tensile testing were also obtained using the TEM under the same settings. The TEM specimens for post-tensile testing were obtained from the section that only suffered uniform deformation for the fractured tensile samples.
2. Experimental
The results of the Vickers microhardness testing performed on each set of specimens upon annealing are presented in Fig. 1. The microhardness values increase steadily with the number of rolling cycles till that of the 6 cycles sample, beyond which the Vickers hardness showed an apparent drop, see the 8 cycles sample measurement. This suggests the 8 cycles specimen consists of a more stable microstructure compared to that of the 6 cycles specimen. The trends of hardness-temperature relation for all specimens are similar, i.e. the hardness value decreases gradually until the temperature of 200 ◦ C. This corresponds to the recovery process in the materials. The hardness decreases at a much higher rate, i.e. at a steeper slope on the curves, from 200 to 300 ◦ C. It levels out past 300 ◦ C, which suggests the onset of rapid grain growth is between 200 and 300 ◦ C. This observation is refined by introducing an annealing at 250 ◦ C to further narrow down the changes within this range. As will be shown in the tensile properties, annealing at 250 ◦ C yielded results still obeying the trend shown by the hardness curves in Fig. 1 for the 200–300 ◦ C range.
Sheets of fully annealed (at 400 ◦ C for 1 h) 1100 series aluminum with a thickness of 1 mm were used as the starting material for the ARB process. The surfaces of the sheets where the interface would be formed were degreased with acetone and wire brushed to enhance the bonding of the two sheets. The two sheets were then stacked one on top of the other and secured to each other by wires and then rolled to 50% reduction at room temperature. This cycle was repeated up to 8 cycles to obtain a von Mises equivalent strain of 6.4. Samples were taken after 1 cycle, 4 cycles, 6 cycles, and 8 cycles of ARB for testing. Subsequent annealing was done in an air furnace at temperatures ranging from 100 to 600 ◦ C for 1.8 ks. The samples were air cooled after annealing. DSC was also performed on the sample after 4 cycles of ARB at 40 K/min in Argon atmosphere. Vickers microhardness values were taken to show the general trend of mechanical behavior upon annealing. Mean Vickers microhardness values were obtained from the average of 10 indentations at random locations on the long transverse face of the samples. A load of 98.07 mN for 10 s was used to obtain the microhardness values. The microhardness results of samples after 6 cycles of ARB showed a similar trend to that of 4 and 8 cycles samples, hence 6 cycles samples will be omitted in further testing. Tensile samples were machined to specification of ASTM standards B-557M subsize specimen scaled down to 40%. The gage length of the samples was 10 mm with a width of 2 mm. The tensile samples were oriented parallel to the rolling direction. The tensile tests were conducted at a strain rate of 7.6 × 10−4 s−1 in ambient temperature with a MTS 810 frame. The microstructure of the specimens under various conditions was viewed under a Hitachi H-800 transmission electron microscope (TEM) with an accelerating voltage of 200 kV. The thin foil samples were prepared by twin-jet electropolishing with an electrolyte of 33% HNO3 and 67% methanol. Both the planar and transverse sections of the samples were viewed. Grain size measurements were performed on transverse view of the samples. The grain size both parallel to and normal to the rolling direction was measured. At least 200 grains from
3. Results 3.1. Mechanical behavior through the use of microhardness testing
Fig. 1. Mean Vickers microhardness of specimens after different number of ARB cycles and conventional aluminum upon annealing.
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Fig. 2. Mean 0.2% offset strength of specimens after different number of ARB cycles upon annealing compared to conventional 1100 series aluminum obtained from ASM Handbook, 8th ed., vol. 2.
3.2. Mechanical behavior through the use of tensile testings Relationships of the 0.2% offset yield strength vs. annealing temperature for specimen after 1, 4, 6, and 8 cycles of ARB are shown in Fig. 2. The individual tensile curves for 1, 4, and 8 cycles specimen in as rolled condition, upon annealing at 250 and 400 ◦ C for 1.8 ks are shown in Fig. 3. The yield strengths measured in the present study, shown in Fig. 2, were noticeably higher than that reported in Ref. [14], although the trend is similar. The increment in strength after 1 cycle of ARB is greatest from about 40 MPa for the annealed AA1100 to about 170 MPa for the 1 cycle sample, and such strength increment decreases after each subsequent ARB cycles. The elongation also showed a similar trend of a steep reduction after 1 cycle and no significant changes after subsequent ARB cycles. Tsuji et al. [16] showed that the elongation of ARB materials could be recovered after the grain sizes reached 1 m. Such a grain size, i.e. 1 m grain size, corresponds to the microstructures after the annealing temperature of just past 200 ◦ C for all specimens in the present study. As shown in Fig. 2, similar drop in yield strength and tensile strength, and decreases in elongation as in the work of Tsuji et al. [16] was observed. It must be pointed out that the 8 cycles specimen under the as-rolled condition has a lower 0.2% offset strength than that of the 6 cycles specimen, compare Fig. 2. It is interesting that the difference in yield strength among the 1 cycle, 4 cycles, and 8 cycles of ARB decreased considerably after annealing at 250 ◦ C for 1.8 ks from the as-rolled condition. The three values are 91, 99, 104 MPa, respectively. As can be seen, although the gaps between the 1, 4, and 8 cycles specimen have been shown to be closer to each other, the difference still exhibits a proportional relationship with the number of ARB cycles it had been through. When the annealing temperature was higher than 300 ◦ C, the yield strengths of the 1 cycle and 4 cycles samples reached similar values as that for the con-
Fig. 3. Tensile curves of specimens after different number of cycles of ARB (a) in the as-rolled condition, (b) upon annealing at 250 ◦ C, and (c) upon annealing at 400 ◦ C.
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ventional annealed 1100 sample. Whereas the 8 cycles sample showed a slightly different trend and the yield strength did not reach the conventional until past 400 ◦ C. Specifically, upon annealing at 250 ◦ C, the tensile curves exhibit a yield point phenomenon for the 4 cycles and 8 cycles specimen (Fig. 3(b)). The short plateau on the curves for the 4 and 8 cycles samples suggest the appearance and propagation of L¨uder’s bands through out the gage length. The same, but less obvious, phenomenon is observed for 4 cycles and 8 cycles specimens annealed at 300 ◦ C as well, although not shown in the present report. The L¨uder strain for 300 ◦ C annealed samples is noticeably lower than that of the 250 ◦ C samples. Upon annealing at 400 ◦ C (Fig. 3(c)), the tensile curves of the specimen that had undergone the ARB process exhibit curves comparable to that of conventional 1100 aluminum. This is also true for specimens annealed at 300 ◦ C excluding the yield point phenomenon mentioned above. Although the 0.2% yield strength for 1 and 4 cycles samples have reached that of conventional 1100 aluminum, the proportionality of UTS values and the number of ARB cycles is still observable in the tensile curves. 3.3. Microstructure evaluation under TEM The micrographs of planar view as well as transverse view for the 4 and 8 cycles samples are shown in Figs. 4 and 5, respectively. The as-rolled specimens after 4 and 8 cycles show similar structure. The dislocation density within the interior of the grain is low in both samples, whereas in the vicinity of the grain boundaries dislocation debris are the main feature. This microstructure differs from that of the as-rolled specimens after only 1 cycle, not shown in this report due to limited space, where the dislocation density within grains are high. The amount of dislocations and dislocation debris decreased after annealing at 200 ◦ C (Figs. 4(b) and 5(b)). The amount of dislocations and dislocation debris decreases significantly after annealing at 250 ◦ C. Upon annealing at higher temperature, i.e. 300 and 400 ◦ C, the microstructure is comparable to that of conventional 1100 aluminum under the same condition. Transverse view micrographs show grain growth in the direction normal to the rolling direction as well. The number of grains between layer interfaces decrease from approximately 5 to 2, from as rolled condition to upon annealing at 250 ◦ C, respectively. The grains have become mostly equiaxed upon annealing at 250 ◦ C and above. Results of the grain size measurement along both directions, parallel and normal to the rolling direction, are presented in Fig. 6. The grain size parallel to the rolling direction obtained in this study is consistent with that presented in Ref. [14]. The curves show continual increase in grain size with increase of the annealing temperature in both directions. This suggests that this material exhibits continuous growth and does not have an apparent conventional recrystallization stage. Such an observation is confirmed by DSC test results for the sample after 4 cycles of ARB. The DSC results, presented in Fig. 7, show no
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apparent exothermic peaks along the temperature range which suggests the process of nucleation is not present. A more interesting phenomenon is that the grain size parallel and normal to the rolling directions increases at similar rate. As a result, the absolute difference between the grain sizes in both directions remains similar as the grain grows as a whole. This process continues until the absolute difference becomes insignificant relative to the grain size of the material, at which point the grain can be perceived as equiaxed. The state of equiaxial grains is achieved upon annealing at 250 ◦ C. After annealing at 300 and 400 ◦ C, it was determined that the grain size for 4 cycles samples reached that of conventional aluminum. The 8 cycles samples upon 300 or 400 ◦ C annealing, however, showed a grain size of approximately ∼2.77–3.01 m and upon annealing at 600 ◦ C, the grain size increased to that of conventional aluminum. Specifically, the cross section of 4 cycles and 8 cycles also revealed some discontinuous segregations of smaller grains located at the layer interfaces (Figs. 4(b) and (c) and 5(b) and (d)). These segregations contain elongated grains that are smaller than that of the surrounding area (approximately 12% of the matrix’s grain size for specimen annealed at 200 and 250 ◦ C) even after annealing at 250–400 ◦ C. The appearance of such segregation does not occur along the entire length of all layer interfaces. The size of these segregations is seemingly random. Upon annealing, the segregated grains do not seem to interact with the matrix of the material (i.e. grain boundaries do not seem to migrate out from or into the segregation). This would imply the existence of some sort of barrier between the segregation and the matrix grains, most likely consisting of oxide always found on the surface of aluminum alloys. 3.4. Post-deformation microstructure under TEM The microstructures for conventional 1100 series aluminum before and after deformation are presented in Fig. 8. The microstructures of the 4 and 8 cycles specimens of as-rolled condition, and 250 and 400 ◦ C annealed conditions after tensile deformation are shown in Figs. 9 and 10, respectively. The TEM thin foils of the deformed samples were taken only from the uniform-deformation portion of the gauge so that the approximate strain levels of these samples could be estimated. The microstructure for the as-rolled sample shows an increase in the amount of dislocations near the grain boundaries. Within the grains, however, the dislocation density did not seem to have increased significantly. Upon deformation of annealed ARB samples, the amount of dislocations within the grains seems to increase with the annealing temperature, the higher the annealing temperature, the bigger the dislocation density in the grain interior under the same applied strain level. Naturally, for the same annealing condition, the higher strain accommodated in the tensile test, the bigger the dislocation density in the grain interior. When the applied strain reached certain high levels depending on the annealing temperature, dislocation cell formation is observed. These features are comparable to those of conventional 1100 series aluminum after uniform deformation.
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Fig. 4. Planar and the corresponding transverse micrographs of specimen after 4 cycles of ARB in (a) as rolled condition, and upon annealing at (b) 200 ◦ C, (c) 250 ◦ C, (d) 400 ◦ C for 1.8 ks. Note: arrow denotes the interface segregation in (b)–(d).
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Fig. 5. Planar and the corresponding transverse micrographs of specimen after 8 cycles of ARB in (a) as rolled condition, and upon annealing at (b) 200 ◦ C, (c) 250 ◦ C, (d) 400 ◦ C for 1.8 ks. Note: arrow denotes the interface segregation in (b)–(d).
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Fig. 7. DSC curve of 1100 aluminum after 4 cycles of ARB compared to DSC curve of conventional 1100 aluminum. DSC scanning rate was set at 40 K/min.
Fig. 6. Grain size parallel and perpendicular to rolling direction of specimens after (a) 4 cycles, (b) 8 cycles of ARB upon annealing at different temperature for 1.8 ks.
4. Discussion 4.1. The introduction of layer interfaces and interlayer segregates The grain size of 1 and 4 cycles specimens upon annealing at 300 ◦ C or higher has been qualitatively shown to be similar to conventional 1100 aluminum. Differently, specimens
of 8 cycles have grain size smaller than that of conventional 1100 aluminum. It is extremely important to point out that the grain size (∼3 m) for the 8 cycles samples measured in the present work is congruous to the calculated separation (3.90 m) between layer interfaces for 8 cycles of ARB starting with a 1 mm sheet. This observation is strong evidence that the introduction of layer boundaries within the bulk does have an effect on the grain growth behavior. Cao et al. [11] has suggested that the alumina particles from the interfaces provided Zener drag on the grain boundaries retarding grain growth. This effect is also observed in the microhardness results. Upon annealing at 600 ◦ C, the grain size appears to be the same as in the conventional aluminum samples. This implies that the materials with layer thickness small enough to produce noticeable retarding effect to the grain growth experience two stages of grain growth. In the first stage, the grains grow until the grain boundaries are close enough to the layer interface and the interaction with the particles would occur. In the latter stage, it is likely that the grain boundaries may overcome the particles’ drag and grow in all directions upon high driving force at high temperatures, leading to the disappearance of some or all of the layer interfaces. Aside from the appearance of oxide segregations or oxide dispersoids, discontinuous segregations of small grains were also found along the layer interfaces. The randomness in size
Fig. 8. Microstructural comparison of conventional 1100 Aluminum specimen before and after uniform deformation.
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Fig. 9. TEM micrographs of 4 cycles ARB samples after maximum uniform tensile deformation (i.e. to UTS point). (a) Sample in the as-rolled condition, (b) sample in 250 ◦ C annealed condition, and (c) sample in 400 ◦ C annealed condition. Note: compare the non-deformed micrographs in Fig. 4.
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Fig. 10. TEM micrographs of 8 cycles ARB samples after maximum uniform tensile deformation (i.e. to UTS point). (a) Sample in the as-rolled condition, (b) sample in 250 ◦ C annealed condition, and (c) sample in 400 ◦ C annealed condition. Note: compare the micrographs of the undeformed specimens in Fig. 5.
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of these segregations seems to be related to the amount of strain accumulated in the local region. The randomness arises from the non-uniform surface deformation accumulated during the wire-brushing step. The non-uniform surface roughness also encourages non-uniform deformation during the following roll bonding processes. The number of these segregation regions and the size of the grains within these segregates seemingly increase with the annealing temperature. The dislocations introduced presumably either dynamically recover during the deformation process promoting the formation of nano-sized grains or they experience recovery during the annealing process. In both cases, they would then grow into the segregates near the interfaces. The intermixed oxide dispersoids might have also retarded the growth of these grains within the segregates, whereas the grains in the matrix of the material did not show similar trend. As can be seen in Fig. 5(c) and (d) (for the transverse micrograph after 250 and 400 ◦ C annealing, respectively), the grains in the segregate seem to maintain still the elongated shape even if the grains in the matrix have become equiaxed. This would suggest the possibility of polygonization rather than simply recovery. Similar phenomenon has also been observed in Ref. [20] with dynamic recrystallization in ARB copper sheets. Another possibility also exists that these grains maintained its shape due the oxide dispersoids acting as restraint for grain boundary migration. Upon annealing at higher temperature, the segregates are mostly engulfed by one or two grains. The segregates do not seem to interact with the bulk until annealing is performed at higher temperatures such that the layer interfaces begin to disappear. The grains within the segregation seem to have difficulty surpassing the circumscribed oxide to grow into the bulk until higher temperature annealing is performed. The introduction of these discontinuous segregates is also proven to be detrimental to the mechanical properties of the sheet as a bulk. Xing et al. [9] has shown that the layer interfaces are the weakest locations in the material, delamination of the lastly formed bond was observed and seems to have cause failure of the material. Furthermore, the appearance of these discontinuous segregates at the already weak interfaces, which are less ductile than the surrounding matrix, will further debilitate the mechanical properties. A large strain compatibility mismatch would appear, when an external strain is applied, between the small grains within the segregation and the large grains outside the segregation. The reduced ability to accommodate the strain compatibly between these grains would cause the formation of voids along the boundaries shared by the large and small grains. Indubitably the oxide that is present between the segregation and the bulk would further weaken the ability to accommodate strain between the segregations and the bulk. The voids that form along the weak interfaces would most likely cause early failure in the material. 4.2. Dynamic microstructural relaxation at higher ARB cycles The microstructure of the 1100 aluminum after 8 cycles of ARB is concluded to be at a lower energy state than that after
6 cycles. This is evident from the results of both the microhardness and tensile testing. It should also be noted that the 8 cycles microstructure is seemingly more stable when annealing is performed at low temperatures. This is also shown in the microhardness and tensile results in Figs. 1 and 2, respectively. Another feature observed is the interface segregation found in the as-rolled 8 cycles sample. Such segregations are also found in 6 cycles specimens, although much less frequently. In this case, the accumulated strain was dynamically recovered during the ARB process either due to the mechanical work or the adiabatic heating that Huang et al. [21] reported to occurred during the ARB process. 4.3. Pertaining to the observed yield point phenomenon The yield point phenomenon seen in the 4 and 8 cycles samples after annealing at 250 ◦ C, and to a lesser extent in the 300 and 400 ◦ C annealed samples, is usually seen in BCC materials such as low carbon steel but it is rarely seen in fcc materials such as aluminum. However, a similar phenomenon as that detected in the present study had been reported by Wyrzykowski and Grabski [22] as well as by Kwieci´nski and Wyrzykowski [23] for pure aluminum with similar grain size. Lloyd et al. [24] and Lloyd and Morris [25] have also reported such phenomena in aluminum alloys. Although no specific mechanisms have been confirmed, they had indeed suggested possible explanations for this yield point phenomenon. Lloyd et al. [24] argued that the occurrence of L¨uder band is due to the lack of mobile dislocations and the lack of dislocation sources thereof. The possibility that the grain boundaries could have influence the yield point behavior was suggested in Ref. [24] as well. Kwieci´nski and Wyrzykowski [23] have also accredited the influence of special low diffusivity grain boundaries on the yield point behavior. The annihilation of dislocations in grain boundaries during annealing would then change the properties of the grain boundaries [23]. It has also been suggested that grain size may play a role in the existence of the yield point phenomenon. Lloyd et al. [24] suggested that an increase in grain size resulted in a lower L¨uder strain; this phenomenon is also seen when comparing the results for the 250, 300 and 400 ◦ C annealed samples in the present study. The L¨uder strain for 250 ◦ C is greater than that of 300 and 400 ◦ C annealed, whereas the similar grain size between 300 and 400 ◦ C yielded a similar L¨uder strain. The possibility of all materials experiencing this yield point behavior at sufficiently small grain size was mentioned in Ref. [22]. In the present investigation, the phenomenon does not appear in the as-rolled samples, nor in samples annealed at lower than 250 ◦ C. Hence the new finding of this study is that the yield point phenomenon seemingly only occurs within a certain range of grain size rather than just below a certain grain size. When the grain size is too large, the L¨uder strain would be much too small to be noticed and the yield point phenomenon becomes insignificant. Thus, the present yield point phenomenon appears to be caused by the combination of the grain size effect and the grain boundary structures influence.
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Fig. 11. Comparison of Hall–Petch plots based on (a) grain size parallel to rolling direction, (b) grain size perpendicular to rolling direction, (c) data for pure polycrystalline aluminum presented by Armstrong et al. [24], (d) data for ARB 1100 aluminum presented by Tsuji et al. [15], and (e) based on this study taking into account texture consideration.
4.4. Microstructural evolution and its implication on the mechanical behavior Tsuji et al. [16] have shown that the 1100 aluminum material of 6 cycles ARB samples follows a Hall–Petch relationship. For the present study, the 0.2% offset yield strength for 4, 6, and 8 cycles are presented in Fig. 11 in a Hall–Petch plot, and the corresponding Hall–Petch constants are also listed. It should be noted that the obtained Hall–Petch constants and friction stress values depend on the grain size measured in specific direction, i.e. parallel or normal to the rolling direction, curve (a) and curve (b), respectively, in Fig. 11. The values for both scenarios in Fig. 11 are found to be quite different from that reported by Armstrong [26] for pure aluminum; However these values are comparable to those reported by Tsuji et al. [16]. In the present work, the strong texture developed during the ARB process would affect the mechanical behavior and it should be somehow integrated into the Hall–Petch analysis. The inadequacy of ignoring the texture effect can be seen in the resulting negative o value using the grain size parallel to the rolling direction, curve (a) in Fig. 11. It has been previously shown in Ref. [11] that the 1100 series aluminum after ARB retained similar texture before and after annealing. In that report, the textures are of -fiber type, and the predominate S ({1 2 3}6 3 4) and copper ({1 1 2}1 1 1) texture cover approximately 80% of the material. With the knowledge of the crystallographic direction of loading in the present case (i.e. parallel to the rolling direction), the paths of the dislocation glide could be estimated using the crystallographic relationship between the loading direction and the glide directions. The glide paths in the textured grains would travel through thickness of the grains at specific angles.
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This suggests that the dislocation movement would be limited by the thickness of the grains rather than the planar grain size. Therefore, the grain size measured in the glide directions should be used to adjust for the texture effect. With the above understanding, the average of the grain size in the glide directions weighted by the orientations and its volume fraction as reported in [11] for 1100 aluminum after 7 cycles of ARB is then used to calculate d, the distance of dislocation movement. The resulting Hall–Petch plot seems to be more realistic compared to the plot without texture consideration, see curve (e) in Fig. 11. The small non-equiaxed grains in the as-rolled samples seem to have altered the mechanical behavior of the material under static uniaxial loading. The TEM micrographs of this material condition after uniform plastic deformation show the lack of dislocations within the grains. This is most likely due to the constraining effect of the small grain size to dislocation movement. As a matter of fact, it is our belief that, in the present case, the mean free path of dislocation–dislocation interactions is longer than the distance between the dislocations and the nearby grain boundaries. Once plastic flow occurs, dislocations would encounter the grain boundaries first before any interactions between themselves could occur; thus formation of dislocation networks of any sort is minimized. It is also a possibility that the stress field produced by the presence of the grain boundaries is interacting with the nearby dislocations within the grains and attracted them toward the grain boundaries. As the annealing temperature increases and the amount of strain that the bulk is capable of accommodating increase accordingly, the dislocation density within the grains seems to increase. The increase in dislocation density in the 400 ◦ C caused an increase in thickness of the member of the dislocation networks. 4.5. The role of microstructure on strain hardening behavior The effect of small grain size on the strain hardening behavior can be clearly seen in Fig. 12 where the relationship of the strain hardening rate within the uniform deformation regime with the true strain level is shown. By comparing the initial strain hardening rate of the as-rolled specimens of different cycles as well as by comparing the annealed ARB samples with the conventional samples, the grain size effect can be clearly seen. The initial strain hardening rate for the as-rolled specimens, in all cases, shows a much larger value than those of the conventional and 400 ◦ C annealed specimens. This high hardening rate at the onset of plastic deformation in the ARB samples is most likely due to the short distance for dislocations to travel before encountering obstacles, i.e. grain boundaries in this case. The small grain size limits also the distance between dislocations, which in turn would most likely contribute to the increase in resistance for further deformation. The effect of grain size can be further seen when comparing the initial strain hardening rate of the asrolled specimens that have gone through different number of ARB cycles. Through the comparison of the initial strain hardening rate, as seen in Fig. 12, and the associated decrease in grain size with each cycle of ARB, it should be clear that the initial strain hardening rate is inversely proportional with the decrease
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the TEM micrographs of the post-deformation samples show a low dislocation density within the grains with some dislocation debris around the grain boundaries. This further conforms to the mechanical behavior within the short uniform deformation regime in the as-rolled condition. Upon annealing at 400 ◦ C the specimen elongation has reached similar values to conventional aluminum. However the initial strain hardening rate for the specimen that has gone through the ARB process is still considerably higher, as shown in Fig. 13. The reason behind this phenomenon still has not been confirmed; however it is obviously related to the processing history of the material. The possibilities include but are not limited to texture, mobile dislocation density and availability of dislocation sources. Further studies would have to be done on this subject matter. 5. Conclusions
Fig. 12. Comparison of strain hardening rate in the as-rolled condition after different number of cycles of ARB.
in grain size. It must be noted again that the grain size difference between 8 cycle as-rolled samples and 4 cycles as-rolled samples is relatively small, yet an increase in initial strain hardening rate is observed; Thus the initial strain hardening rate is shown to be also directly proportional to the accumulated strain. The lack of ability to accommodate strain would lead to problems associated with strain incompatibility with neighboring grains, thus leading to the development of voids and cracks at grain boundaries. The development of cracks and voids would then limit the elongation to a relatively low value, which corresponds well with the 2% elongation of the ARBed samples compared to the 35% elongation in the conventional aluminum. Evidently,
Fig. 13. Comparison of strain hardening rate in the as-rolled condition after different number of cycles of ARB upon annealing at 400 ◦ C.
The mechanical behavior of commercial purity 1100 series aluminum after different number of accumulative roll-bonding cycles, followed by subsequent annealing, was studied and related to the microstructure of the material. The following conclusions have been reached: (1) Through the use of proper heat treatment conditions (such as 250 ◦ C annealed in this study), it was observed that it is possible to increase the ductility significantly without sacrificing large amount of strength of the ARB 1100 aluminum samples. (2) Rapid grain growth was detected only when the annealing temperature was higher than 150 ◦ C. The grain growth rates below 200 ◦ C show similar trends in both the rolling direction and in the normal direction. (3) The oxide rolled into the material at the interfaces has a tendency of retarding grain boundary migration through the interface and it then retards grain growth. Such conclusion is clearly evidenced by the grain size data for the 8 cycles samples upon annealing at 300 and 400 ◦ C, and further supported by the mechanical testing results. (4) The strain introduced during the surface treatment prior to each pass and the nature of the bonding mechanism creates a layer of high strain around the interface. This layer of high strain would then lead to recovery or possibly polygonization upon annealing to form discontinuous segregation containing smaller grains within. There were also evidence of embedded oxide with and surrounding these segregates. (5) Similar discontinuous segregations of small grains are found in the as-rolled 8 cycles specimens. In this case, the energy needed to introduce segregation is supplied in the form of mechanical work or adiabatic heating during the ARB process. (6) Yield point phenomenon is detected in the ARB specimens after 250 ◦ C annealing, even if both the grain refinement process and the resultant grain morphology in this case are very different from that in the previous reports whereby yield point phenomenon was observed. This study has further
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demonstrated that such yield point phenomenon occurred only in a specific range of grain size. (7) The strong texture accumulated during ARB was taken into account in to the Hall–Petch relationship and was found to have better represented the strength and grain size relationship of ARB materials. Acknowledgement This project was funded by Natural Science and Engineering Research Council of Canada (NSERC). References [1] J. Mao, S.B. Kang, J.O. Park, J. Mater. Process. Technol. 159 (2005) 314–320. [2] R.Z. Valiev, D.A. Salimonenko, N.K. Tsenev, P.B. Berbon, T.G. Langdon, Scripta Mater. 37 (1997) 1945–1950. [3] S. Lee, P.B. Berbon, M. Furukawa, Z. Horita, M. Nemoto, N.K. Tsenev, R.Z. Valiev, T.G. Langdon, Mater. Sci. Eng. A A272 (1999) 63–72. [4] N. Tsuji, K. Shiotsuki, H. Utsunomiya, Y. Saito, Mater. Sci. Forum 304–306 (1999) 73–78. [5] Y. Saito, N. Tsuji, H. Utsunomiya, T. Sakai, R.G. Hong, Scripta Mater. 39 (1998) 1221–1227. [6] Y. Saito, H. Utsunomiya, N. Tsuji, T. Sakai, Acta Mater. 47 (1999) 579–583. [7] N. Tsuji, Y. Saito, S.-H. Lee, Y. Minamino, Adv. Eng. Mater. 5 (2003) 338–344.
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[8] Z.P. Xing, S.B. Kang, H.W. Kim, J. Mater. Sci. 39 (2004) 1259–1265. [9] Z.P. Xing, S.B. Kang, H.W. Kim, Metall. Mater. Trans. A 33 (2002) 1521–1530. [10] S.-H. Lee, C.H. Lee, C.Y. Lim, Mater. Sci. Forum 449–452 (2004) 161–164. [11] W.Q. Cao, Q. Liu, A. Godfrey, N. Hansen, Mater. Sci. Forum 408–412 (2002) 721–726. [12] C.P. Heason, P.B. Prangnell, Mater. Sci. Forum 408–412 (2002) 733–738. [13] X. Huang, N. Tsuji, N. Hansen, Y. Minamino, Mater. Sci. Eng. A 340 (2003) 265–271. [14] Y.-S. Kim, T.-O. Lee, D.H. Shin, Mater. Sci. Forum 449–452 (2004) 625–628. [15] N. Tsuji, Y. Ito, Y. Saito, Y. Minamino, Scripta Mater. 47 (2002) 893–899. [16] N. Tsuji, S. Okuno, T. Matsuura, Y. Koizumi, Y. Minamino, Mater. Sci. Forum 426–432 (2003) 2667–2672. [17] H.W. H¨oppel, J. May, M. G¨oken, Adv. Eng. Mater. 6 (2004) 219–222. [18] N. Tsuji, Y. Ito, H. Nakashim, F. Yoshida, Y. Minamino, Mater. Sci. Forum 396–402 (2002) 423–428. [19] R. Ueji, X. Huang, N. Hansen, N. Tsuji, Y. Minamino, Mater. Sci. Forum 426–432 (2003) 405–410. [20] B.L. Li, N. Tsuji, N. Kamikawa, Mater. Sci. Eng. A A423 (2006) 331–342. [21] X. Huang, N. Hansen, N. Tsuji, Science 312 (2006) 249–251. [22] J.W. Wyrzykowski, M.W. Grabski, Mater. Sci. Eng. 56 (1982) 197–200. [23] J. Kwieci´nski, J.W. Wyrzykowski, Acta Metall. Mater. 41 (1993) 3089–3095. [24] D.J. Lloyd, S.A. Court, K.M. Gatenby, Mater. Sci. Technol. 13 (1997) 660–666. [25] D.J. Lloyd, L.R. Morris, Acta Metall. 25 (1977) 857–861. [26] R.W. Armstrong, Adv. Mater. Res. 4 (1970) 101–146.
Mechanical behavior and microstructural evolution upon annealing of the accumulative roll-bonding (ARB) processed Al alloy 1100 Charles Kwan a , Zhirui Wang a,∗ , Suk-Bong Kang b a
b
Department of Materials Science and Engineering, University of Toronto, 184 College Street, Toronto, Ont., Canada M2J 1L2 Materials Engineering Department, Korea Institute of Machinery and Materials, 66 Sangnam, Kyungnam, Changwon 641-010, Republic of Korea Received 17 April 2007; received in revised form 27 June 2007; accepted 3 July 2007
Abstract The results of mechanical testing on ultra-fine grained aluminum processed by accumulative roll-bonding (ARB) were analyzed with TEM observation and in accordance of the microstructural evolution upon annealing. It was found that rapid grain growth, with the corresponding decrease in strength, did not occur until the annealing temperature of 200 ◦ C or higher. The oxide rolled into the material near the bonding interfaces was seen to act as an obstruction for grain boundary migration across said interfaces. More interestingly, the strain near the interface due to the surface preparation technique used during ARB was found to form discontinuous segregates consisting of smaller grains formed during annealing or even ARB processing of higher number of cycles. Such phenomenon is attributed to recovery or polygonization due to the strain incurred. This study has also demonstrated that yield point phenomenon may be observed in a commercially pure fcc metal when the grain size is within a certain range. © 2007 Elsevier B.V. All rights reserved. Keywords: Accumulative roll-bonding; Aluminum; Mechanical behavior; Annealing; Microstructure evolution; Strain hardening
1. Introduction Materials with ultra-fine grain (UFG) structure (having a mean grain size between 100 nm and 1 m) have been gathering the interest of scientists for more than 50 years and especially in the past 20 years. Not only do the UFGed materials display increase in strength through Hall–Petch strengthening, but also these materials have high strain rate super plasticity [1–3] and low temperature super plasticity [4]. It has been shown that UFGed structure can be obtained through severe plastic deformation, methods such as torsion straining and equal channel angular pressing (ECAP). Most of these available processes lack the potential for a cost efficient scale up for the production of bulk UFGed material in large quantity. Saito et al. [5,6] have successfully created a method to form bulk UFGed structure through a processing technique called accumulative roll-bonding (ARB). The ARB process is essentially a process of multiple rolling, whereby after each roll to a 50% reduction in the sheet thickness, the sheet is cut into halves, surface treated and stacked up before
∗
Corresponding author. Tel.: +1 416 978 4412; fax: +1 416 978 4415. E-mail address: [email protected] (Z. Wang).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.07.022
it is rolled again [5,6]. By stacking the two halves between each 50% reduction roll, the thickness of the sheet remains unchanged throughout any given number of cycles. This process has been successfully applied to make UFGed sheets of aluminum and aluminum alloys [7–15], and IF steel [15,16]. Hall–Petch strengthening is based on the increase in the increased grain boundary volume fraction acting as obstacles for dislocation glide. Through the use of the ARB process, the tensile strength of 1100 aluminum reached 270 MPa after 6 cycles of ARB, corresponding to an increase of approximately 3 times from the 90 MPa tensile strength of the base material [14,15,17]. On the contrary, the elongation to failure decreases from approximately 40% to 5% upon even 1 cycle of ARB [14]. Although H¨oppel et al. have shown an increase of elongation to failure as the number of ARB cycles increases [17], the results presented by them strongly support the need of post-ARB annealing in order to recover at least some ductility. The microstructure resulting from the ARB process is shown to be in a pancake-like shaped grains surrounded by well-defined high angle boundaries [8–16]. In recent years, specifically, Huang et al. [13], Tsuji et al. [18], and Ueji et al. [19], showed that the lamellar boundaries in ARB materials consisted mostly of high angle boundaries, whereas the interconnecting
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boundaries between the pancake grains (running vertical through thickness) consisted of a mixture of low and high angle boundaries. Obviously, such high angle boundaries parallel to the rolling direction may bring on new and unique mechanical behavior to these materials. In addition, the ARB process also introduces layer interfaces within the bulk, and for materials that readily form oxide film, such as aluminum, the layer interfaces introduce particles (alumina in this case) into the bulk, which should also alter the mechanical behavior as well as the grain growth behavior. Up to date, there has been a lack of systematic studies on the mechanical behavior of ARB materials, and even less is known about the changes of microstructure as well as mechanical behavior upon annealing of this type of materials. Hence, the present study on the subject matter has been carried out.
three or more distinctive micrographs were measured to obtain the average grain size. Grain size of 1, 4, and 6 cycles samples upon annealing at 300 ◦ C or higher could not be obtained this way, as the grain size under these annealing conditions was too large to contain more than one or two grains in single TEM micrographs. Micrographs after tensile testing were also obtained using the TEM under the same settings. The TEM specimens for post-tensile testing were obtained from the section that only suffered uniform deformation for the fractured tensile samples.
2. Experimental
The results of the Vickers microhardness testing performed on each set of specimens upon annealing are presented in Fig. 1. The microhardness values increase steadily with the number of rolling cycles till that of the 6 cycles sample, beyond which the Vickers hardness showed an apparent drop, see the 8 cycles sample measurement. This suggests the 8 cycles specimen consists of a more stable microstructure compared to that of the 6 cycles specimen. The trends of hardness-temperature relation for all specimens are similar, i.e. the hardness value decreases gradually until the temperature of 200 ◦ C. This corresponds to the recovery process in the materials. The hardness decreases at a much higher rate, i.e. at a steeper slope on the curves, from 200 to 300 ◦ C. It levels out past 300 ◦ C, which suggests the onset of rapid grain growth is between 200 and 300 ◦ C. This observation is refined by introducing an annealing at 250 ◦ C to further narrow down the changes within this range. As will be shown in the tensile properties, annealing at 250 ◦ C yielded results still obeying the trend shown by the hardness curves in Fig. 1 for the 200–300 ◦ C range.
Sheets of fully annealed (at 400 ◦ C for 1 h) 1100 series aluminum with a thickness of 1 mm were used as the starting material for the ARB process. The surfaces of the sheets where the interface would be formed were degreased with acetone and wire brushed to enhance the bonding of the two sheets. The two sheets were then stacked one on top of the other and secured to each other by wires and then rolled to 50% reduction at room temperature. This cycle was repeated up to 8 cycles to obtain a von Mises equivalent strain of 6.4. Samples were taken after 1 cycle, 4 cycles, 6 cycles, and 8 cycles of ARB for testing. Subsequent annealing was done in an air furnace at temperatures ranging from 100 to 600 ◦ C for 1.8 ks. The samples were air cooled after annealing. DSC was also performed on the sample after 4 cycles of ARB at 40 K/min in Argon atmosphere. Vickers microhardness values were taken to show the general trend of mechanical behavior upon annealing. Mean Vickers microhardness values were obtained from the average of 10 indentations at random locations on the long transverse face of the samples. A load of 98.07 mN for 10 s was used to obtain the microhardness values. The microhardness results of samples after 6 cycles of ARB showed a similar trend to that of 4 and 8 cycles samples, hence 6 cycles samples will be omitted in further testing. Tensile samples were machined to specification of ASTM standards B-557M subsize specimen scaled down to 40%. The gage length of the samples was 10 mm with a width of 2 mm. The tensile samples were oriented parallel to the rolling direction. The tensile tests were conducted at a strain rate of 7.6 × 10−4 s−1 in ambient temperature with a MTS 810 frame. The microstructure of the specimens under various conditions was viewed under a Hitachi H-800 transmission electron microscope (TEM) with an accelerating voltage of 200 kV. The thin foil samples were prepared by twin-jet electropolishing with an electrolyte of 33% HNO3 and 67% methanol. Both the planar and transverse sections of the samples were viewed. Grain size measurements were performed on transverse view of the samples. The grain size both parallel to and normal to the rolling direction was measured. At least 200 grains from
3. Results 3.1. Mechanical behavior through the use of microhardness testing
Fig. 1. Mean Vickers microhardness of specimens after different number of ARB cycles and conventional aluminum upon annealing.
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Fig. 2. Mean 0.2% offset strength of specimens after different number of ARB cycles upon annealing compared to conventional 1100 series aluminum obtained from ASM Handbook, 8th ed., vol. 2.
3.2. Mechanical behavior through the use of tensile testings Relationships of the 0.2% offset yield strength vs. annealing temperature for specimen after 1, 4, 6, and 8 cycles of ARB are shown in Fig. 2. The individual tensile curves for 1, 4, and 8 cycles specimen in as rolled condition, upon annealing at 250 and 400 ◦ C for 1.8 ks are shown in Fig. 3. The yield strengths measured in the present study, shown in Fig. 2, were noticeably higher than that reported in Ref. [14], although the trend is similar. The increment in strength after 1 cycle of ARB is greatest from about 40 MPa for the annealed AA1100 to about 170 MPa for the 1 cycle sample, and such strength increment decreases after each subsequent ARB cycles. The elongation also showed a similar trend of a steep reduction after 1 cycle and no significant changes after subsequent ARB cycles. Tsuji et al. [16] showed that the elongation of ARB materials could be recovered after the grain sizes reached 1 m. Such a grain size, i.e. 1 m grain size, corresponds to the microstructures after the annealing temperature of just past 200 ◦ C for all specimens in the present study. As shown in Fig. 2, similar drop in yield strength and tensile strength, and decreases in elongation as in the work of Tsuji et al. [16] was observed. It must be pointed out that the 8 cycles specimen under the as-rolled condition has a lower 0.2% offset strength than that of the 6 cycles specimen, compare Fig. 2. It is interesting that the difference in yield strength among the 1 cycle, 4 cycles, and 8 cycles of ARB decreased considerably after annealing at 250 ◦ C for 1.8 ks from the as-rolled condition. The three values are 91, 99, 104 MPa, respectively. As can be seen, although the gaps between the 1, 4, and 8 cycles specimen have been shown to be closer to each other, the difference still exhibits a proportional relationship with the number of ARB cycles it had been through. When the annealing temperature was higher than 300 ◦ C, the yield strengths of the 1 cycle and 4 cycles samples reached similar values as that for the con-
Fig. 3. Tensile curves of specimens after different number of cycles of ARB (a) in the as-rolled condition, (b) upon annealing at 250 ◦ C, and (c) upon annealing at 400 ◦ C.
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ventional annealed 1100 sample. Whereas the 8 cycles sample showed a slightly different trend and the yield strength did not reach the conventional until past 400 ◦ C. Specifically, upon annealing at 250 ◦ C, the tensile curves exhibit a yield point phenomenon for the 4 cycles and 8 cycles specimen (Fig. 3(b)). The short plateau on the curves for the 4 and 8 cycles samples suggest the appearance and propagation of L¨uder’s bands through out the gage length. The same, but less obvious, phenomenon is observed for 4 cycles and 8 cycles specimens annealed at 300 ◦ C as well, although not shown in the present report. The L¨uder strain for 300 ◦ C annealed samples is noticeably lower than that of the 250 ◦ C samples. Upon annealing at 400 ◦ C (Fig. 3(c)), the tensile curves of the specimen that had undergone the ARB process exhibit curves comparable to that of conventional 1100 aluminum. This is also true for specimens annealed at 300 ◦ C excluding the yield point phenomenon mentioned above. Although the 0.2% yield strength for 1 and 4 cycles samples have reached that of conventional 1100 aluminum, the proportionality of UTS values and the number of ARB cycles is still observable in the tensile curves. 3.3. Microstructure evaluation under TEM The micrographs of planar view as well as transverse view for the 4 and 8 cycles samples are shown in Figs. 4 and 5, respectively. The as-rolled specimens after 4 and 8 cycles show similar structure. The dislocation density within the interior of the grain is low in both samples, whereas in the vicinity of the grain boundaries dislocation debris are the main feature. This microstructure differs from that of the as-rolled specimens after only 1 cycle, not shown in this report due to limited space, where the dislocation density within grains are high. The amount of dislocations and dislocation debris decreased after annealing at 200 ◦ C (Figs. 4(b) and 5(b)). The amount of dislocations and dislocation debris decreases significantly after annealing at 250 ◦ C. Upon annealing at higher temperature, i.e. 300 and 400 ◦ C, the microstructure is comparable to that of conventional 1100 aluminum under the same condition. Transverse view micrographs show grain growth in the direction normal to the rolling direction as well. The number of grains between layer interfaces decrease from approximately 5 to 2, from as rolled condition to upon annealing at 250 ◦ C, respectively. The grains have become mostly equiaxed upon annealing at 250 ◦ C and above. Results of the grain size measurement along both directions, parallel and normal to the rolling direction, are presented in Fig. 6. The grain size parallel to the rolling direction obtained in this study is consistent with that presented in Ref. [14]. The curves show continual increase in grain size with increase of the annealing temperature in both directions. This suggests that this material exhibits continuous growth and does not have an apparent conventional recrystallization stage. Such an observation is confirmed by DSC test results for the sample after 4 cycles of ARB. The DSC results, presented in Fig. 7, show no
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apparent exothermic peaks along the temperature range which suggests the process of nucleation is not present. A more interesting phenomenon is that the grain size parallel and normal to the rolling directions increases at similar rate. As a result, the absolute difference between the grain sizes in both directions remains similar as the grain grows as a whole. This process continues until the absolute difference becomes insignificant relative to the grain size of the material, at which point the grain can be perceived as equiaxed. The state of equiaxial grains is achieved upon annealing at 250 ◦ C. After annealing at 300 and 400 ◦ C, it was determined that the grain size for 4 cycles samples reached that of conventional aluminum. The 8 cycles samples upon 300 or 400 ◦ C annealing, however, showed a grain size of approximately ∼2.77–3.01 m and upon annealing at 600 ◦ C, the grain size increased to that of conventional aluminum. Specifically, the cross section of 4 cycles and 8 cycles also revealed some discontinuous segregations of smaller grains located at the layer interfaces (Figs. 4(b) and (c) and 5(b) and (d)). These segregations contain elongated grains that are smaller than that of the surrounding area (approximately 12% of the matrix’s grain size for specimen annealed at 200 and 250 ◦ C) even after annealing at 250–400 ◦ C. The appearance of such segregation does not occur along the entire length of all layer interfaces. The size of these segregations is seemingly random. Upon annealing, the segregated grains do not seem to interact with the matrix of the material (i.e. grain boundaries do not seem to migrate out from or into the segregation). This would imply the existence of some sort of barrier between the segregation and the matrix grains, most likely consisting of oxide always found on the surface of aluminum alloys. 3.4. Post-deformation microstructure under TEM The microstructures for conventional 1100 series aluminum before and after deformation are presented in Fig. 8. The microstructures of the 4 and 8 cycles specimens of as-rolled condition, and 250 and 400 ◦ C annealed conditions after tensile deformation are shown in Figs. 9 and 10, respectively. The TEM thin foils of the deformed samples were taken only from the uniform-deformation portion of the gauge so that the approximate strain levels of these samples could be estimated. The microstructure for the as-rolled sample shows an increase in the amount of dislocations near the grain boundaries. Within the grains, however, the dislocation density did not seem to have increased significantly. Upon deformation of annealed ARB samples, the amount of dislocations within the grains seems to increase with the annealing temperature, the higher the annealing temperature, the bigger the dislocation density in the grain interior under the same applied strain level. Naturally, for the same annealing condition, the higher strain accommodated in the tensile test, the bigger the dislocation density in the grain interior. When the applied strain reached certain high levels depending on the annealing temperature, dislocation cell formation is observed. These features are comparable to those of conventional 1100 series aluminum after uniform deformation.
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Fig. 4. Planar and the corresponding transverse micrographs of specimen after 4 cycles of ARB in (a) as rolled condition, and upon annealing at (b) 200 ◦ C, (c) 250 ◦ C, (d) 400 ◦ C for 1.8 ks. Note: arrow denotes the interface segregation in (b)–(d).
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Fig. 5. Planar and the corresponding transverse micrographs of specimen after 8 cycles of ARB in (a) as rolled condition, and upon annealing at (b) 200 ◦ C, (c) 250 ◦ C, (d) 400 ◦ C for 1.8 ks. Note: arrow denotes the interface segregation in (b)–(d).
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Fig. 7. DSC curve of 1100 aluminum after 4 cycles of ARB compared to DSC curve of conventional 1100 aluminum. DSC scanning rate was set at 40 K/min.
Fig. 6. Grain size parallel and perpendicular to rolling direction of specimens after (a) 4 cycles, (b) 8 cycles of ARB upon annealing at different temperature for 1.8 ks.
4. Discussion 4.1. The introduction of layer interfaces and interlayer segregates The grain size of 1 and 4 cycles specimens upon annealing at 300 ◦ C or higher has been qualitatively shown to be similar to conventional 1100 aluminum. Differently, specimens
of 8 cycles have grain size smaller than that of conventional 1100 aluminum. It is extremely important to point out that the grain size (∼3 m) for the 8 cycles samples measured in the present work is congruous to the calculated separation (3.90 m) between layer interfaces for 8 cycles of ARB starting with a 1 mm sheet. This observation is strong evidence that the introduction of layer boundaries within the bulk does have an effect on the grain growth behavior. Cao et al. [11] has suggested that the alumina particles from the interfaces provided Zener drag on the grain boundaries retarding grain growth. This effect is also observed in the microhardness results. Upon annealing at 600 ◦ C, the grain size appears to be the same as in the conventional aluminum samples. This implies that the materials with layer thickness small enough to produce noticeable retarding effect to the grain growth experience two stages of grain growth. In the first stage, the grains grow until the grain boundaries are close enough to the layer interface and the interaction with the particles would occur. In the latter stage, it is likely that the grain boundaries may overcome the particles’ drag and grow in all directions upon high driving force at high temperatures, leading to the disappearance of some or all of the layer interfaces. Aside from the appearance of oxide segregations or oxide dispersoids, discontinuous segregations of small grains were also found along the layer interfaces. The randomness in size
Fig. 8. Microstructural comparison of conventional 1100 Aluminum specimen before and after uniform deformation.
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Fig. 9. TEM micrographs of 4 cycles ARB samples after maximum uniform tensile deformation (i.e. to UTS point). (a) Sample in the as-rolled condition, (b) sample in 250 ◦ C annealed condition, and (c) sample in 400 ◦ C annealed condition. Note: compare the non-deformed micrographs in Fig. 4.
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Fig. 10. TEM micrographs of 8 cycles ARB samples after maximum uniform tensile deformation (i.e. to UTS point). (a) Sample in the as-rolled condition, (b) sample in 250 ◦ C annealed condition, and (c) sample in 400 ◦ C annealed condition. Note: compare the micrographs of the undeformed specimens in Fig. 5.
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of these segregations seems to be related to the amount of strain accumulated in the local region. The randomness arises from the non-uniform surface deformation accumulated during the wire-brushing step. The non-uniform surface roughness also encourages non-uniform deformation during the following roll bonding processes. The number of these segregation regions and the size of the grains within these segregates seemingly increase with the annealing temperature. The dislocations introduced presumably either dynamically recover during the deformation process promoting the formation of nano-sized grains or they experience recovery during the annealing process. In both cases, they would then grow into the segregates near the interfaces. The intermixed oxide dispersoids might have also retarded the growth of these grains within the segregates, whereas the grains in the matrix of the material did not show similar trend. As can be seen in Fig. 5(c) and (d) (for the transverse micrograph after 250 and 400 ◦ C annealing, respectively), the grains in the segregate seem to maintain still the elongated shape even if the grains in the matrix have become equiaxed. This would suggest the possibility of polygonization rather than simply recovery. Similar phenomenon has also been observed in Ref. [20] with dynamic recrystallization in ARB copper sheets. Another possibility also exists that these grains maintained its shape due the oxide dispersoids acting as restraint for grain boundary migration. Upon annealing at higher temperature, the segregates are mostly engulfed by one or two grains. The segregates do not seem to interact with the bulk until annealing is performed at higher temperatures such that the layer interfaces begin to disappear. The grains within the segregation seem to have difficulty surpassing the circumscribed oxide to grow into the bulk until higher temperature annealing is performed. The introduction of these discontinuous segregates is also proven to be detrimental to the mechanical properties of the sheet as a bulk. Xing et al. [9] has shown that the layer interfaces are the weakest locations in the material, delamination of the lastly formed bond was observed and seems to have cause failure of the material. Furthermore, the appearance of these discontinuous segregates at the already weak interfaces, which are less ductile than the surrounding matrix, will further debilitate the mechanical properties. A large strain compatibility mismatch would appear, when an external strain is applied, between the small grains within the segregation and the large grains outside the segregation. The reduced ability to accommodate the strain compatibly between these grains would cause the formation of voids along the boundaries shared by the large and small grains. Indubitably the oxide that is present between the segregation and the bulk would further weaken the ability to accommodate strain between the segregations and the bulk. The voids that form along the weak interfaces would most likely cause early failure in the material. 4.2. Dynamic microstructural relaxation at higher ARB cycles The microstructure of the 1100 aluminum after 8 cycles of ARB is concluded to be at a lower energy state than that after
6 cycles. This is evident from the results of both the microhardness and tensile testing. It should also be noted that the 8 cycles microstructure is seemingly more stable when annealing is performed at low temperatures. This is also shown in the microhardness and tensile results in Figs. 1 and 2, respectively. Another feature observed is the interface segregation found in the as-rolled 8 cycles sample. Such segregations are also found in 6 cycles specimens, although much less frequently. In this case, the accumulated strain was dynamically recovered during the ARB process either due to the mechanical work or the adiabatic heating that Huang et al. [21] reported to occurred during the ARB process. 4.3. Pertaining to the observed yield point phenomenon The yield point phenomenon seen in the 4 and 8 cycles samples after annealing at 250 ◦ C, and to a lesser extent in the 300 and 400 ◦ C annealed samples, is usually seen in BCC materials such as low carbon steel but it is rarely seen in fcc materials such as aluminum. However, a similar phenomenon as that detected in the present study had been reported by Wyrzykowski and Grabski [22] as well as by Kwieci´nski and Wyrzykowski [23] for pure aluminum with similar grain size. Lloyd et al. [24] and Lloyd and Morris [25] have also reported such phenomena in aluminum alloys. Although no specific mechanisms have been confirmed, they had indeed suggested possible explanations for this yield point phenomenon. Lloyd et al. [24] argued that the occurrence of L¨uder band is due to the lack of mobile dislocations and the lack of dislocation sources thereof. The possibility that the grain boundaries could have influence the yield point behavior was suggested in Ref. [24] as well. Kwieci´nski and Wyrzykowski [23] have also accredited the influence of special low diffusivity grain boundaries on the yield point behavior. The annihilation of dislocations in grain boundaries during annealing would then change the properties of the grain boundaries [23]. It has also been suggested that grain size may play a role in the existence of the yield point phenomenon. Lloyd et al. [24] suggested that an increase in grain size resulted in a lower L¨uder strain; this phenomenon is also seen when comparing the results for the 250, 300 and 400 ◦ C annealed samples in the present study. The L¨uder strain for 250 ◦ C is greater than that of 300 and 400 ◦ C annealed, whereas the similar grain size between 300 and 400 ◦ C yielded a similar L¨uder strain. The possibility of all materials experiencing this yield point behavior at sufficiently small grain size was mentioned in Ref. [22]. In the present investigation, the phenomenon does not appear in the as-rolled samples, nor in samples annealed at lower than 250 ◦ C. Hence the new finding of this study is that the yield point phenomenon seemingly only occurs within a certain range of grain size rather than just below a certain grain size. When the grain size is too large, the L¨uder strain would be much too small to be noticed and the yield point phenomenon becomes insignificant. Thus, the present yield point phenomenon appears to be caused by the combination of the grain size effect and the grain boundary structures influence.
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Fig. 11. Comparison of Hall–Petch plots based on (a) grain size parallel to rolling direction, (b) grain size perpendicular to rolling direction, (c) data for pure polycrystalline aluminum presented by Armstrong et al. [24], (d) data for ARB 1100 aluminum presented by Tsuji et al. [15], and (e) based on this study taking into account texture consideration.
4.4. Microstructural evolution and its implication on the mechanical behavior Tsuji et al. [16] have shown that the 1100 aluminum material of 6 cycles ARB samples follows a Hall–Petch relationship. For the present study, the 0.2% offset yield strength for 4, 6, and 8 cycles are presented in Fig. 11 in a Hall–Petch plot, and the corresponding Hall–Petch constants are also listed. It should be noted that the obtained Hall–Petch constants and friction stress values depend on the grain size measured in specific direction, i.e. parallel or normal to the rolling direction, curve (a) and curve (b), respectively, in Fig. 11. The values for both scenarios in Fig. 11 are found to be quite different from that reported by Armstrong [26] for pure aluminum; However these values are comparable to those reported by Tsuji et al. [16]. In the present work, the strong texture developed during the ARB process would affect the mechanical behavior and it should be somehow integrated into the Hall–Petch analysis. The inadequacy of ignoring the texture effect can be seen in the resulting negative o value using the grain size parallel to the rolling direction, curve (a) in Fig. 11. It has been previously shown in Ref. [11] that the 1100 series aluminum after ARB retained similar texture before and after annealing. In that report, the textures are of -fiber type, and the predominate S ({1 2 3}6 3 4) and copper ({1 1 2}1 1 1) texture cover approximately 80% of the material. With the knowledge of the crystallographic direction of loading in the present case (i.e. parallel to the rolling direction), the paths of the dislocation glide could be estimated using the crystallographic relationship between the loading direction and the glide directions. The glide paths in the textured grains would travel through thickness of the grains at specific angles.
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This suggests that the dislocation movement would be limited by the thickness of the grains rather than the planar grain size. Therefore, the grain size measured in the glide directions should be used to adjust for the texture effect. With the above understanding, the average of the grain size in the glide directions weighted by the orientations and its volume fraction as reported in [11] for 1100 aluminum after 7 cycles of ARB is then used to calculate d, the distance of dislocation movement. The resulting Hall–Petch plot seems to be more realistic compared to the plot without texture consideration, see curve (e) in Fig. 11. The small non-equiaxed grains in the as-rolled samples seem to have altered the mechanical behavior of the material under static uniaxial loading. The TEM micrographs of this material condition after uniform plastic deformation show the lack of dislocations within the grains. This is most likely due to the constraining effect of the small grain size to dislocation movement. As a matter of fact, it is our belief that, in the present case, the mean free path of dislocation–dislocation interactions is longer than the distance between the dislocations and the nearby grain boundaries. Once plastic flow occurs, dislocations would encounter the grain boundaries first before any interactions between themselves could occur; thus formation of dislocation networks of any sort is minimized. It is also a possibility that the stress field produced by the presence of the grain boundaries is interacting with the nearby dislocations within the grains and attracted them toward the grain boundaries. As the annealing temperature increases and the amount of strain that the bulk is capable of accommodating increase accordingly, the dislocation density within the grains seems to increase. The increase in dislocation density in the 400 ◦ C caused an increase in thickness of the member of the dislocation networks. 4.5. The role of microstructure on strain hardening behavior The effect of small grain size on the strain hardening behavior can be clearly seen in Fig. 12 where the relationship of the strain hardening rate within the uniform deformation regime with the true strain level is shown. By comparing the initial strain hardening rate of the as-rolled specimens of different cycles as well as by comparing the annealed ARB samples with the conventional samples, the grain size effect can be clearly seen. The initial strain hardening rate for the as-rolled specimens, in all cases, shows a much larger value than those of the conventional and 400 ◦ C annealed specimens. This high hardening rate at the onset of plastic deformation in the ARB samples is most likely due to the short distance for dislocations to travel before encountering obstacles, i.e. grain boundaries in this case. The small grain size limits also the distance between dislocations, which in turn would most likely contribute to the increase in resistance for further deformation. The effect of grain size can be further seen when comparing the initial strain hardening rate of the asrolled specimens that have gone through different number of ARB cycles. Through the comparison of the initial strain hardening rate, as seen in Fig. 12, and the associated decrease in grain size with each cycle of ARB, it should be clear that the initial strain hardening rate is inversely proportional with the decrease
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the TEM micrographs of the post-deformation samples show a low dislocation density within the grains with some dislocation debris around the grain boundaries. This further conforms to the mechanical behavior within the short uniform deformation regime in the as-rolled condition. Upon annealing at 400 ◦ C the specimen elongation has reached similar values to conventional aluminum. However the initial strain hardening rate for the specimen that has gone through the ARB process is still considerably higher, as shown in Fig. 13. The reason behind this phenomenon still has not been confirmed; however it is obviously related to the processing history of the material. The possibilities include but are not limited to texture, mobile dislocation density and availability of dislocation sources. Further studies would have to be done on this subject matter. 5. Conclusions
Fig. 12. Comparison of strain hardening rate in the as-rolled condition after different number of cycles of ARB.
in grain size. It must be noted again that the grain size difference between 8 cycle as-rolled samples and 4 cycles as-rolled samples is relatively small, yet an increase in initial strain hardening rate is observed; Thus the initial strain hardening rate is shown to be also directly proportional to the accumulated strain. The lack of ability to accommodate strain would lead to problems associated with strain incompatibility with neighboring grains, thus leading to the development of voids and cracks at grain boundaries. The development of cracks and voids would then limit the elongation to a relatively low value, which corresponds well with the 2% elongation of the ARBed samples compared to the 35% elongation in the conventional aluminum. Evidently,
Fig. 13. Comparison of strain hardening rate in the as-rolled condition after different number of cycles of ARB upon annealing at 400 ◦ C.
The mechanical behavior of commercial purity 1100 series aluminum after different number of accumulative roll-bonding cycles, followed by subsequent annealing, was studied and related to the microstructure of the material. The following conclusions have been reached: (1) Through the use of proper heat treatment conditions (such as 250 ◦ C annealed in this study), it was observed that it is possible to increase the ductility significantly without sacrificing large amount of strength of the ARB 1100 aluminum samples. (2) Rapid grain growth was detected only when the annealing temperature was higher than 150 ◦ C. The grain growth rates below 200 ◦ C show similar trends in both the rolling direction and in the normal direction. (3) The oxide rolled into the material at the interfaces has a tendency of retarding grain boundary migration through the interface and it then retards grain growth. Such conclusion is clearly evidenced by the grain size data for the 8 cycles samples upon annealing at 300 and 400 ◦ C, and further supported by the mechanical testing results. (4) The strain introduced during the surface treatment prior to each pass and the nature of the bonding mechanism creates a layer of high strain around the interface. This layer of high strain would then lead to recovery or possibly polygonization upon annealing to form discontinuous segregation containing smaller grains within. There were also evidence of embedded oxide with and surrounding these segregates. (5) Similar discontinuous segregations of small grains are found in the as-rolled 8 cycles specimens. In this case, the energy needed to introduce segregation is supplied in the form of mechanical work or adiabatic heating during the ARB process. (6) Yield point phenomenon is detected in the ARB specimens after 250 ◦ C annealing, even if both the grain refinement process and the resultant grain morphology in this case are very different from that in the previous reports whereby yield point phenomenon was observed. This study has further
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demonstrated that such yield point phenomenon occurred only in a specific range of grain size. (7) The strong texture accumulated during ARB was taken into account in to the Hall–Petch relationship and was found to have better represented the strength and grain size relationship of ARB materials. Acknowledgement This project was funded by Natural Science and Engineering Research Council of Canada (NSERC). References [1] J. Mao, S.B. Kang, J.O. Park, J. Mater. Process. Technol. 159 (2005) 314–320. [2] R.Z. Valiev, D.A. Salimonenko, N.K. Tsenev, P.B. Berbon, T.G. Langdon, Scripta Mater. 37 (1997) 1945–1950. [3] S. Lee, P.B. Berbon, M. Furukawa, Z. Horita, M. Nemoto, N.K. Tsenev, R.Z. Valiev, T.G. Langdon, Mater. Sci. Eng. A A272 (1999) 63–72. [4] N. Tsuji, K. Shiotsuki, H. Utsunomiya, Y. Saito, Mater. Sci. Forum 304–306 (1999) 73–78. [5] Y. Saito, N. Tsuji, H. Utsunomiya, T. Sakai, R.G. Hong, Scripta Mater. 39 (1998) 1221–1227. [6] Y. Saito, H. Utsunomiya, N. Tsuji, T. Sakai, Acta Mater. 47 (1999) 579–583. [7] N. Tsuji, Y. Saito, S.-H. Lee, Y. Minamino, Adv. Eng. Mater. 5 (2003) 338–344.
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