10-03-80A-4.qxd
27/4/04
17:24
Page 1735
Achieving Enhanced Ductility in a Dilute Magnesium Alloy through Severe Plastic Deformation KIYOSHI MATSUBARA, YUICHI MIYAHARA, ZENJI HORITA, and TERENCE G. LANGDON Experiments were conducted to evaluate the utility of a new processing procedure developed for Mg-based alloys in which samples are subjected to a two-step processing route of extrusion followed by equal-channel angular pressing (designated as EX-ECAP). The experiments were conducted using a Mg-0.6 wt pct Zr alloy and, for comparison purposes, samples of pure Mg. It is shown that the potential for successfully using ECAP increases in both materials when adopting the EX-ECAP procedure. For the Mg-Zr alloy, the use of EX-ECAP produces a grain size of ⬃1.4 m when the pressing is undertaken at 573 K. By contrast, using EX-ECAP with pure Mg at 573 K produces a grain size of ⬃26 m. Tensile testing of the Mg-Zr alloy at 523 and 573 K after processing by EX-ECAP revealed the occurrence of significantly enhanced ductilities with maximum elongations of ⬃300 to 400 pct.
I.
INTRODUCTION
THE low density and good machinability of magnesium alloys makes them attractive for a wide range of structural applications in transportation and, especially, in the automotive field.[1,2] These alloys also exhibit excellent damping capacities, so that they are especially attractive for use in applications where it is necessary to dampen external vibrations, as in helicopters and satellites. These beneficial features have led to the prediction by Friedrich and Schumann[3] of a “new age of magnesium.” Despite the inherent potential advantages of magnesiumbased alloys, their hexagonal crystal structure provides only a limited number of active slip systems, giving low levels of ductility and, consequently, difficulties in forming complex components. It is well established in recent work that the ductilities of metallic alloys may be significantly enhanced, often to the superplastic range, by processing the alloys through the introduction of severe plastic deformation (SPD).[4] Conventional methods of SPD processing include equal-channel angular pressing (ECAP),[5,6,7] where a sample is pressed repetitively through a die confined within a channel bent through an angle at, or close to, 90 deg, and high-pressure torsion (HPT),[8,9,10] where a disk is subjected to a high pressure and concurrent torsional straining. The procedures of ECAP and HPT are both effective in producing substantial grain refinement in fcc metals such as pure aluminum,[11,12] pure nickel,[9,10] and a range of aluminum alloys,[13–16] but, in practice, ECAP appears to be the more useful technique because it utilizes fairly large samples, it can be scaled up relatively easily to produce large KIYOSHI MATSUBARA and YUICHI MIYAHARA, Graduate Students, and ZENJI HORITA, Professor, are with the Department of Materials Science and Engineering, Faculty of Engineering, Kyushu University, Fukuoka 812-8581, Japan. TERENCE G. LANGDON, Professor, is with the Departments of Aerospace & Mechanical Engineering and Materials Science, University of Southern California, Los Angeles, CA 90089-1453. Contact e-mail:
[email protected] Manuscript submitted February 18, 2003. METALLURGICAL AND MATERIALS TRANSACTIONS A
bulk material,[17] and high strains may be imposed through the use of a multipass die[18] or by using a rotary die[19] or a side-extrusion facility.[20] In addition, it has been demonstrated that the principles of ECAP processing may be incorporated into conventional rolling using the conshearing[21] or continuous confined strip-shearing (C2S2)[22–25] processes. In contrast to the work with fcc metals, there has been less success when using ECAP processing with magnesium and Mg-based alloys. Although there are several reports of the application of ECAP to a range of Mg alloys,[26–30] these investigations generally relate to the use of fairly complex alloy systems where there are widespread dispersions of intermetallic precipitates. By contrast, the processing by ECAP of as-cast pure Mg and dilute Mg solid-solution alloys has been generally unsuccessful, and ultrafine-grained microstructures have not been achieved. For example, pure Mg was subjected to ECAP through two passes at 673 K to give a measured as-pressed grain size of ⬎100 m, and there was a measured grain size of ⬃17 m in an Mg-0.9 pct Al solid-solution alloy after ECAP through two passes at 473 K,[31] where this and all subsequent alloy compositions are given in weight percent. It is apparent, therefore, that both of these grain sizes are outside of the range of conventional ultrafine-grained materials, so that, although the experiments on pure Mg and the Mg-0.9 pct Al alloy demonstrated an improvement in the mechanical properties of these materials at room temperature, the as-pressed grain sizes were too large to give enhanced or superplastic ductilities when testing at elevated temperatures. The present investigation was motivated by recent evidence that very significant grain refinement may be introduced in a Mg-0.6 pct Zr alloy through the use of a two-step processing procedure in which the alloy is first extruded and then subjected to ECAP processing.[32] This two-step process of extrusion and subsequent ECAP has been given the acronym EX-ECAP, and recent experiments have shown that, for the same conditions of temperature and strain rate, this process is capable of increasing the ductility of a Mg-9 pct Al alloy from ⬃30 pct in the as-cast condition to ⬃110 pct after casting plus extrusion and up to ⬃840 pct after casting and processing by EX-ECAP.[33] VOLUME 35A, JUNE 2004—1735
10-03-80A-4.qxd
27/4/04
17:24
Page 1736
The basic principles underlying the EX-ECAP process are threefold. First, early experiments showed that it was possible to use conventional extrusion in order to achieve significant grain refinement in commercial-purity Mg.[34] Second, it is well known that the application of extrusion to Mg-based alloys is often effective in promoting superplastic characteristics in these materials,[35] primarily because the extrusion step produces some limited grain refinement, in addition to a strong texture where the basal planes become reasonably aligned.[36,37] Third, by applying ECAP to samples in the extruded condition, it is reasonable to anticipate that there may be an additional grain refinement and, thus, an ability to attain even higher superplastic elongations. Preliminary experiments on a Mg-0.6 pct Zr alloy showed that this two-step EX-ECAP procedure was effective in reducing the grain size to ⬃1 m through only a single pass of ECAP at 573 K, and the aspressed alloy then exhibited superplastic characteristics when testing in tension at 573 K.[32] The earlier exploratory work on the Mg-0.6 pct Zr alloy was designed only to evaluate the potential for using an EX-ECAP processing route to achieve an ultrafine microstructure and good tensile ductilities in a dilute magnesium-based alloy. It was recognized, however, that a detailed investigation was required in order to more fully characterize the potential of this processing procedure. Accordingly, the present investigation was initiated with four specific objectives. The first was to evaluate the feasibility of using the EX-ECAP procedure on the Mg-0.6 pct Zr alloy with the ECAP conducted at different temperatures and through different numbers of passes, so that varying amounts of strain are introduced in each pressing operation. The second was to examine the thermal stability of any ultrafine-grained microstructure introduced into the alloy. The third objective was to evaluate the potential for achieving enhanced ductility or superplastic elongations in the Mg-0.6 pct Zr alloy after processing by EX-ECAP. The fourth was to make a direct comparison between the microstructures achieved in the Mg-0.6 pct Zr alloy using EX-ECAP with those obtained when using the same processing route with pure Mg.
II.
EXPERIMENTAL MATERIALS AND PROCEDURES
The material used in this investigation was from a different batch, but of the same nominal composition, as the Mg-0.6 wt pct Zr alloy documented in an earlier report.[32] As previously, this alloy was selected specifically because it has been demonstrated that this represents the optimum composition in order to achieve a very high damping capacity.[38] The alloy was produced through a conventional melting and casting process, and it was supplied in the form of a cylindrical ingot having a diameter of 60 mm and a length of ⬃100 mm. A chemical analysis of the cast alloy gave a zirconium content of ⬃0.7 wt pct with an addition of ⬃0.01 wt pct Al but with no other measurable impurities. Since the maximum solubility of Zr in cast Mg is ⬃0.3 wt pct from room temperature to ⬃573 K, it is apparent that the alloy contained a small quantity of an ␣-Zr phase.[39] Inspection showed the measured grain size in the cast alloy to be ⬃70 m. For comparison purposes, some additional experiments were conducted using pure Mg, which was also 1736—VOLUME 35A, JUNE 2004
prepared by casting into an ingot of 60 mm in diameter. The grain size of the cast Mg was estimated to be ⬃1.4 mm, and analysis revealed the presence of 0.02 wt pct Al as the only measurable impurity. The melting temperature of the Mg-Zr alloy was determined using differential scanning calorimetry (DSC). A small sample of the alloy, with a weight of ⬃50 mg, was heated in an argon atmosphere together with a Mo standard within a MacScience DSC3300S apparatus, using a heating rate of 10 K min⫺1. The incipient melting temperature of the alloy was recorded as 920 K. To evaluate the effect of incorporating extrusion into the processing procedure, parts of the Mg-Zr alloy and the pure Mg were extruded at a speed of 5 mm s⫺1 and a temperature of 623 K into rods having a diameter of 10 mm, where this extrusion condition corresponds to a reduction ratio of 36:1. These materials are, henceforth, designated as the “cast ⫹ extruded” condition. For ECAP processing, samples were prepared with diameters of 10 mm and lengths of 60 mm, and the pressing was conducted using a solid die with an angle between the channels of ⌽ ⫽ 90 deg and an outer arc of curvature of ⌿ ⬇ 20 deg at the point where the two parts of the channel intersect. It can be shown from first principles that these internal angles give an imposed strain of ⬃1 on each passage of the sample through the die,[40] and experiments have demonstrated that an arc of curvature of ⌿ ⬇ 20 deg has no significant influence on the subsequent homogeneity of the microstructure.[41] When samples were pressed repetitively through the die, each sample was rotated by 90 deg in the same sense between each pass in the processing procedure designated as route BC.[42,43] All pressings were conducted at elevated temperatures, with the samples held at temperature within the die for ⬃10 minutes prior to the first pass and then, when repetitive pressings were undertaken, they were held at temperature within the die for ⬃1 minute before pressing. Attempts were made to press the Mg-Zr alloy for up to four passes using temperatures in the range from 473 to 573 K. Preliminary experiments showed that higher temperatures were needed for the pressing of pure Mg, and, for this material, the pressing temperatures ranged from 573 to 673 K. Some samples were subjected to ECAP after casting but without an intermediate extrusion; these samples are designated as “cast ⫹ equal-channel angular pressed.” Other samples were prepared by casting, extruded to the diameter required for ECAP, and then subjected to ECAP at selected temperatures through up to a maximum of four passes: the latter samples represent the EX-ECAP process, and these samples are designated as “cast ⫹ extruded ⫹ equal-channel angular pressed.” The thermal stability of the microstructure was examined for samples of the Mg-Zr alloy subjected to EX-ECAP processing by slicing samples with thicknesses of ⬃0.4 mm perpendicular to the longitudinal axes after EX-ECAP and then sealing these disks in glass tubes in an argon environment and annealing for 1 hour at selected temperatures up to a maximum of 673 K. The microstructures of the Mg-Zr alloy and the pure Mg were subject to extensive examination using combinations of optical microscopy (OM), transmission electron microscopy (TEM), and electron-probe microanalysis (EPMA). For METALLURGICAL AND MATERIALS TRANSACTIONS A
10-03-80A-4.qxd
27/4/04
17:24
Page 1737
inspection by OM, samples were cut parallel to the solidification direction after casting or parallel to the extrusion direction in the cast ⫹ extruded or cast ⫹ extruded ⫹ equalchannel angular pressed conditions, and they were ground mechanically to a mirrorlike surface using abrasive papers and alumina powders. These surfaces were then etched using a solution of 1 pct HNO3, 24 pct C2H5OH, and 75 pct H2O, with an immersion time of ⬃15 to 20 seconds. For TEM, the extruded or pressed rods were sliced perpendicular to their longitudinal axes to thicknesses of ⬃0.4 mm, cut and ground to disks with diameters of 3 mm and thicknesses of ⬃0.15 mm, and then thinned to perforation at room temperature in a twin-jet polishing facility using a solution of 15 pct HClO4, 15 pct C3H8O3, and 70 pct C2H5OH. Following perforation, they were ion milled in an argon environment. The TEM observations were conducted using an Hitachi H-8100 electron microscope operating at 200 kV, and selected-area electron diffraction (SAED) patterns were recorded from regions having diameters of 6.2 m. The microanalysis was undertaken using a Shimadzu EPMA-1600 facility operating at 15 kV with a probe current of ⬃10 nA, and X-ray maps were collected over square areas of 80 ⫻ 80 m2. The tensile properties of the Mg-Zr alloy were evaluated by machining tensile specimens with gage lengths of 5 mm and cross-sectional dimensions of 2 ⫻ 3 mm2. These specimens were prepared with their tensile axes cut parallel to the longitudinal axes after extrusion or EX-ECAP and parallel to the solidification direction in the cast condition. Tensile tests were conducted at 523 and 573 K using initial strain rates in the range from 3.3 ⫻ 10⫺4 to 3.3 ⫻ 10⫺1 s⫺1 and with a testing machine operating at a constant rate of crosshead displacement. All tensile tests were taken to failure to provide information on the total elongations. The values of the strain-rate sensitivity (m) were determined for these two testing temperatures, where m is defined as ⭸ ln s [1] # ⭸ ln # where is the flow stress and is the imposed strain rate.[44] The stress-strain curves were examined to determine the values of the flow stresses at a strain of 0.1 for each testing condition, and these data were then plotted logarithmically as flow stress against strain rate to provide estimates of the values for m. m⫽
III.
EXPERIMENTAL RESULTS
A. Microstructures after Casting and Extrusion Figure 1 shows optical micrographs of the pure Mg (on the left-hand side) and the Mg-0.6 pct Zr alloy (on the right-hand side) in the as-cast and cast ⫹ extruded conditions, respectively: the extrusion direction is horizontal for the samples subjected to extrusion. After casting, many twins are visible in the pure Mg, but extrusion leads to an array of equiaxed grains indicating that recrystallization probably occurs during extrusion at 623 K. Measurements showed that the aver age grain size was reduced from ⬃1.4 mm after casting to ⬃55 m in the cast ⫹ extruded condition. Inspection shows also that the cast Mg-Zr alloy contains second-phase particles, and examination by EPMA demonstrated these particles to METALLURGICAL AND MATERIALS TRANSACTIONS A
be rich in Zr. The presence of these Zr-rich particles is effective in retaining a grain size of only ⬃70 m in the cast condition, and the grain size is further reduced to ⬃11 m in the cast ⫹ extruded condition, where there is again an array of reasonably equiaxed grains suggesting the occurrence of recrystallization. As is evident from Figure 1, dark bands are also visible in the alloy in the cast ⫹ extruded condition, with the bands lying essentially parallel to the extrusion direction: it was shown using EPMA that these bands are also rich in Zr, and it is reasonable to conclude they are instrumental in retaining the grain size at this very low level. B. Nature of the Microstructures after ECAP A detailed overview of the ECAP experiments is given in Table I for both pure Mg and the Mg-Zr alloy: this tabulation distinguishes between samples pressed successfully without any cracking (denoted by X), samples pressed but showing visible macroscopic surface cracking (denoted by A), and samples which broke during the pressing (denoted by B), where the other possible pressing conditions were not attempted. This tabulation confirms the advantage of introducing the extrusion step between casting and ECAP. Thus, the Mg-Zr alloy was successfully pressed through four passes at 573 K in the EX-ECAP condition, but the alloy broke on the second pass when pressed at this temperature in the cast condition without any extrusion. Similarly, the pure Mg was successfully pressed through two passes without cracking at 573 and 623 K using the EXECAP procedure, whereas, when pressing the cast material at these two temperatures, the samples broke in the first and second pass, respectively. The additional strains that may be imposed through ECAP in the cast ⫹ extruded samples is attributed to the significant grain refinement occurring during the extrusion step. It is apparent also that the introduction of 0.6 pct Zr leads to a significant increase in the pressing capability by comparison with pure Mg, and this is also associated with the greater grain refinement that may be achieved in the dilute alloy. Microstructural information for the Mg-Zr alloy is shown in Figure 2, where X-ray maps are presented for the cast, extruded, and EX-ECAP condition, where the latter corresponds to a sample pressed under the optimum condition of four passes at a temperature of 573 K. These maps were constructed using, for the same field of view for any condition, the Mg K␣ (upper row) and Zr L␣ (lower row) lines in the EPMA, and the gradations in color from blue through green to yellow and red depict the increasing local enrichments of the selected elements. Thus, it is apparent that there are Zrrich regions where Mg is depleted in all three conditions, but these Zr-rich regions are well defined and reasonably rounded in cross section in the cast condition and they are well-defined and elongated in the extrusion direction after the extrusion step, but they become diffuse and have ill-defined boundaries after the subsequent ECAP. The microstructure of pure Mg after EX-ECAP is shown in Figure 3, where the sample was pressed for two passes at 573 K. There is an equiaxed array of grains in this sample, with an average size of ⬃26 m. Thus, EX-ECAP refines the grain size in pure Mg, but it is not capable, at least at 573 K, of reducing the grain size to the submicrometer level. Twinning is also evident in this microstructure, but the twins VOLUME 35A, JUNE 2004—1737
10-03-80A-4.qxd
27/4/04
17:24
Page 1738
Fig. 1—Optical micrographs of pure Mg (on left) and the Mg-0.6 pct Zr alloy (on right) in the as-cast and cast ⫹ extruded conditions: the extrusion direction is indicated.
Table I.
Material
Summary of the Effect of Applying ECAP to Pure Mg and the Mg-0.6 pct Zr Alloy in the Cast and the Cast ⫹ Extruded Conditions
Number of Passes
Pure Mg Mg-0.6 pct Zr
1 2 3 1 2 3 4
Cast
Cast ⫹ Extruded
ECAP at T (K)
ECAP at T (K)
573
623
673
B
X B
X X B
X B
473
B
523
A B
573
623
X X B X X X X
X X B
X ⫽ successful pressing without any visible cracking; A ⫽ evidence for macroscopic surface cracking after ECAP; and B ⫽ specimen broken during ECAP.
are less clearly defined than in the as-cast structure shown in Figure 1. By contrast, the microstructures of the Mg-Zr alloy after EX-ECAP were more representative of those achieved in the ECAP of fcc metals. 1738—VOLUME 35A, JUNE 2004
Figure 4 shows examples of the microstructures in the alloy after EX-ECAP with the pressing conducted at 573 K through one, two, and four passes, respectively, together with the appropriate SAED patterns. After one pass, the boundaries METALLURGICAL AND MATERIALS TRANSACTIONS A
10-03-80A-4.qxd
27/4/04
17:24
Page 1739
Fig. 2—EPMA for the Mg-0.6 pct Zr alloy showing the three conditions of as-cast, cast ⫹ extruded and cast ⫹ extruded–equal-channel angular pressed: the upper and lower rows show the characteristics of Mg K␣ and Zr L␣, respectively.
Fig. 3—Microstructure of pure Mg after extrusion and ECAP through two passes at 573 K.
tend to be poorly defined, and there is a very high dislocation density: microstructures with similar characteristics were reported earlier for various fcc metals after SPD processing, METALLURGICAL AND MATERIALS TRANSACTIONS A
and these types of boundaries are generally considered to be in high-energy, nonequilibrium configurations.[14,45–47] The microstructure is inhomogeneous after a single pass, with some areas showing arrays of high-angle boundaries, as demonstrated by the SAED pattern taken at point (a), and other areas showing arrays of subgrains with low-angle boundaries, as demonstrated by the pattern taken at point (b). The diffuse and ill-defined nature of the microstructure after one pass makes it impossible to provide a quantitative estimate of the relevant proportions of the regions containing predominantly high- and low-angle boundaries, but, nevertheless, the overall impression was that both regions were present throughout the sample. After two passes, there are well-defined equiaxed grains in many regions, but also some areas of poorly defined boundaries with high dislocation densities. The measured grain size in this condition was ⬃1.3 m, and it is evident from the SAED pattern in Figure 4 that these grain boundaries have predominantly high angles of misorientation. A well-defined microstructure of equiaxed grains was visible throughout most of the material after four passes, and there were relatively few dislocations within the grains. The grain size in this condition was ⬃1.4 m, and the SAED pattern in Figure 4 suggests the presence of predominantly high-angle boundaries. In practice, a grain size of ⬃1.4 m is slightly VOLUME 35A, JUNE 2004—1739
10-03-80A-4.qxd
27/4/04
17:24
Page 1740
Fig. 4—Microstructures in the Mg-0.6 pct Zr alloy after one, two, and four passes at 573 K, plus appropriate SAED patterns: the two SAED patterns for one pass relate to the areas marked (a) and (b).
outside of the range of conventional ultrafine-grained materials, where the grain sizes are usually defined as ⬍1 m. C. Thermal Stability of the Ultrafine-Grained Microstructure The preceding observations show that the EX-ECAP process is effective in reducing the grain size to close to ⬃1 m in the Mg-0.6 pct Zr alloy, but any attempts to achieve high tensile ductilities at elevated temperatures requires also that these very fine grains are reasonably stable at high temperatures. Figure 5 shows results obtained from annealing samples for 1 hour at selected temperatures up to a maximum of 673 K: these results were obtained with samples prepared by EX-ECAP including four passes at 573 K. It is apparent from these measurements that the fine-grained structure is reasonably stable up to 573 K, but some limited grain growth occurs at higher temperatures. At 623 K, it was observed that fine grains having sizes of ⬍2 m covered an estimated ⬃30 pct of the sample, and these smaller grains tended to be retained along longitudinal stringers lying parallel to the extrusion and pressing directions, as illustrated in Figure 6 where an arrow delineates some of the smaller grains present after annealing at 623 K. A comparison with the EPMA Zr mapping in Figure 2 suggests that these smaller grains are retained preferentially in those regions that are 1740—VOLUME 35A, JUNE 2004
Fig. 5—Grain size vs temperature for samples of the Mg-Zr alloy processed by extrusion and ECAP for four passes at 573 K and then annealed for 1 h at selected temperatures.
Zr-rich, thereby confirming that the presence of the Zr-rich particles prevents significant grain growth. These observations suggest, therefore, that it would be beneficial to achieve METALLURGICAL AND MATERIALS TRANSACTIONS A
10-03-80A-4.qxd
27/4/04
17:24
Page 1741
Fig. 6—Microstructure of the Mg-0.6 pct Zr alloy after preparing by EXECAP including four passes at 573 K and then annealing for 1 h at 623 K: the arrow points to smaller grains contained in longitudinal stringers lying horizontal and parallel to the extrusion and pressing directions.
Fig. 7—Elongation to failure vs strain rate for the Mg-0.6 pct Zr alloy at 573 K in four different processing conditions: in the as-received cast condition, in the cast ⫹ extruded condition, in the cast ⫹ extruded ⫹ equalchannel angular pressed condition and in the cast ⫹ equal-channel angular pressed condition without any intermediate extrusion.
a homogeneous dispersion of Zr-rich particles throughout the material, since this would assist in attaining a uniform distribution of ultrafine grains in the processed alloy. D. Mechanical Properties at Elevated Temperatures Samples were tested in tension under four different microstructural conditions: in the as-received cast condition; in the cast ⫹ extruded condition; in the cast ⫹ equal-channel angular pressed condition, where the ECAP was conducted for one pass at 573 K without any intermediate extrusion; and in the cast ⫹ extruded ⫹ equal-channel angular pressed condition where, in order to provide consistent data with the cast ⫹ equal-channel angular pressed specimen, the samples METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 8—Elongation to failure vs strain rate for the Mg-0.6 pct Zr alloy at 573 K for cast ⫹ extruded samples in an unpressed condition and for samples pressed at 573 K through one, two, and four passes.
were pressed through only one pass at 573 K. All samples were pulled to failure at 573 K, and the results are shown in Figure 7. These results are quite similar to those reported in the earlier investigation using a Mg-0.6 pct Zr alloy, but where the alloy was obtained from a different batch.[32] The overall similarity between these two batches of alloys confirms the general reproducibility of the data. It is evident from Figure 7 that the cast alloy exhibits only limited ductility, but the elongation to failure is increased to a small extent by applying ECAP, is increased to a greater extent if the cast alloy is extruded, and is increased even further when the material is processed using both extrusion and ECAP in the EX-ECAP procedure. Thus, the imposition of only a single pass of ECAP after extrusion has a very significant effect on the elongation to failure, despite the fact that the microstructure after one pass is ill-defined and contains a large number of dislocations within the grains. These results confirm, therefore, the importance of using a two-step processing route when working with Mg-based alloys. The significance of the number of passes on the measured elongations to failure at 573 K is shown in Figure 8, where results are presented for extruded samples in an unpressed condition plus for extruded samples taken through one, two, and four passes of ECAP at 573 K. In general, it appears that the total elongation tends to increase slightly with increasing numbers of passes, and this trend is especially evident at the faster strain rates. This increase is attributed, as shown in Figure 4, to the more homogeneous array of equiaxed grains present in the alloy after the larger number of passes. Taking the measured flow stress for each sample at a strain () of 0.1, Figure 9 shows a logarithmic plot of flow stress vs strain rate for the same samples recorded in Figure 8. Thus, the flow stress is significantly reduced through the application of ECAP and, since these results were obtained at a high temperature in a region where diffusion-controlled processes are dominant, this trend is consistent with the smaller grain sizes in these samples. The strain-rate sensitivity (m) was recorded as ⬃0.4 at the lower strain rates where the elongations are a maximum for those samples subjected to VOLUME 35A, JUNE 2004—1741
10-03-80A-4.qxd
27/4/04
17:24
Page 1742
Fig. 9—Flow stress at a strain of 0.1 vs strain rate for the samples shown in Fig. 8: the strain-rate sensitivity under conditions of maximum ductility is ⬃0.4.
Fig. 11—Flow stress at a strain of 0.1 vs strain rate for the samples shown in Fig. 10: the strain-rate sensitivity under conditions of maximum ductility is ⬃0.3.
these data, Figure 11 shows a logarithmic plot of flow stress against strain rate, where the flow stress was recorded at a strain of ⫽ 0.1. Under these conditions, the flow stress again decreases with increasing numbers of passes in ECAP, and there is a maximum value of m ⬇ 0.3 for the sample subjected to four passes of ECAP when tested at the two lowest strain rates. IV.
Fig. 10—Elongation to failure vs strain rate for the Mg-0.6 pct Zr alloy at 523 K for unpressed samples and for samples pressed at 573 K through one, two, and four passes.
EX-ECAP through four passes. This value of m is intermediate between the values of m of ⬃0.5 and ⬃0.3 reported for fcc aluminum-based alloys tested after ECAP for conditions where the rate-controlling processes were separately attributed to superplasticity and dislocation glide, respectively.[48,49] Similar results are shown in Figure 10 for the same set of samples, but with the testing conducted at a lower temperature of 523 K. These results match those obtained at 573 K in Figure 8, except that the total elongations tend to be slightly lower at all strain rates. The highest elongation at this temperature is ⬃320 pct when using an initial strain rate of 3.3 ⫻ 10⫺4 s⫺1 after EX-ECAP through four passes. When it is noted that the melting temperature of the alloy was measured as 920 K, it is apparent that these tests relate to a homologeous testing temperature of only 0.57 Tm, where Tm is the absolute melting temperature of the alloy. Using 1742—VOLUME 35A, JUNE 2004
DISCUSSION
These experiments provide important information on the processing of dilute Mg-based alloys through ECAP, and they lead to four significant conclusions. First, the EX-ECAP procedure is remarkably effective in increasing the potential for applying the ECAP process to both pure Mg and a dilute Mg-based alloy. It is difficult to use ECAP processing with pure Mg, but, nevertheless, as shown in Table I, the introduction of an intermediate extrusion increases the viability of the ECAP process so that samples can be pressed through two passes without exhibiting any deleterious cracking. For the dilute Mg-0.6 pct Zr alloy, the use of EX-ECAP increases the viability of the pressing procedure at 573 K from a single pass to four passes. Second, the microstructures are substantially refined when an extrusion step is introduced prior to ECAP. Thus, it is possible to achieve a grain size of ⬃26 m in pure Mg after EX-ECAP, as given in Figure 3, whereas, in the absence of an extrusion step, an earlier report showed the application of ECAP to the cast alloy gave an as-pressed grain size of ⬎100 m.[31] In the Mg-0.6 pct Zr alloy, processing by EX-ECAP produced a grain size of ⬃1.4 m, which is within the range of ultrafine grain sizes generally associated with the processing of fcc metals using severe plastic deformation: for example, experiments have recorded a grain size of ⬃1.3 m in pure aluminum after conventional ECAP.[12] It is apparent from the TEM observations and associated EPMA that the grains are much smaller in the Mg-0.6 pct Zr alloy by comparison with pure Mg METALLURGICAL AND MATERIALS TRANSACTIONS A
10-03-80A-4.qxd
27/4/04
17:24
Page 1743
because of the presence of Zr-rich bands and Zr particles that inhibit grain growth. Third, the ultrafine grain size attained in the Mg-0.6 pct Zr alloy is reasonably stable at high temperatures, at least up to ⬃550 K. As illustrated in Figure 5, a grain size of ⬍10 m is retained in the alloy up to temperatures of at least ⬃600 K. This thermal stability provides an opportunity for achieving enhanced ductilities in tensile tests at elevated temperatures. Fourth, the maximum elongations achieved in the Mg0.6 pct Zr alloy lie typically in the range of ⬃300 to 400 pct, as shown in Figures 8 and 10, with higher elongations recorded at the higher testing temperature of 573 K. These elongations are generally not high by comparison with the very large superplastic elongations recorded in fcc metals after conventional ECAP without an extrusion step: for example, elongations of ⬎2000 pct have been reported during tensile testing of an Al-3 pct Mg-0.2 pct Sc alloy after ECAP.[48] Nevertheless, the present elongations compare very favorably with those reported for other Mg-based alloys prepared using alternative processing procedures. For example, there are reports of maximum elongations of 320 pct at a testing temperature of 773 K in a hot-rolled AZ31 (Mg-3 pct Al-1 pct Zn) alloy where dynamic recrystallization occurred during testing,[50] of 360 pct in a cross-rolled AZ31 alloy subjected to dynamic recrystallization at a low temperature and then tensile testing at 723 K,[51] and of 580 pct at 673 K in a hotrolled AZ61 (Mg-6 pct Al-1 pct Zn) alloy.[52] All of these elongations are comparable to those obtained in the present experiments, but the testing temperatures are consistently 100 °C to 200 °C higher. In the present work, the tensile elongations of ⬎300 pct achieved at 573 K, equivalent to 0.57 Tm, are sufficient for use in superplastic forming operations where the requisite strains are equivalent, typically, to tensile strains of the order of ⬃300 to 400 pct.[53,54] Finally, additional clarification is required with reference to two of the experimental observations. First, the measured strain-rate sensitivities in the Mg-0.6 pct Zr alloy are ⬃0.3 to 0.4, and this appears to be more consistent with dislocation glide as the rate-controlling process rather than superplasticity. In dislocation glide, it is well established that the value of n ⫽ 1/m ⫽ 3.0, where n is the stress exponent. By contrast, true superplasticity is generally associated with a value of n ⬇ 2, equivalent to m ⬇ 0.5.[55] Similar observations and conclusions were reported recently for an aluminum-based 2024 (Al-4.4 pct Cu-1.5 pct Mg-0.6 pct Mn) alloy, where m ⬇ 0.3 and the maximum elongations were also ⬃300 to 500 pct after ECAP processing.[49] It has been shown that dislocation glide can be the rate-controlling process during the high-temperature creep of dilute Mg-based alloys when the grain sizes are large: for example, experiments on a Mg-0.8 pct Al alloy with a grain size of ⬃240 m gave n ⬇ 3.0 at temperatures in the range from 523 to 623 K.[56] However, more work is now needed to clarify the precise nature of the flow mechanism in ultrafine-grained materials exhibiting enhanced ductilities with m ⬇ 0.3 to 0.4. Second, it is intriguing to note that the highest elongation recorded in these experiments was ⬃380 pct at 573 K for the alloy subjected to only a single pass of ECAP, as shown in Figure 8. At first sight, this point appears anomalous by reference to the elongations achieved after processing through two and four passes. However, an even higher elongation METALLURGICAL AND MATERIALS TRANSACTIONS A
of ⬃420 pct was recorded for the same alloy at the same testing temperature and strain rate in an earlier investigation, when the material was also subjected to one pass at 573 K.[32] It is reasonable to speculate that these higher ductilities may be a consequence of an inhibition in grain growth during tensile testing in samples having inhomogeneous microstructures with a preponderance of high-energy, nonequilibrium boundaries, but more experiments are needed to clarify the significance of these exceptionally high elongations at the slowest experimental strain rate. V.
SUMMARY AND CONCLUSIONS
1. Experiments were conducted to examine the feasibility of achieving an ultrafine-grained structure in a Mg-0.6 pct Zr alloy using ECAP. For comparison purposes, additional experiments were also conducted using samples of pure Mg. 2. The potential for successfully using ECAP is markedly increased in both the Mg-Zr alloy and pure Mg by using the EX-ECAP process, in which the material is subjected to extrusion prior to processing by ECAP. When using EX-ECAP, pure Mg may be pressed through two passes without breaking at temperatures of 573 and 623 K, and the Mg-Zr alloy may be pressed through four passes without breaking at 573 K. 3. Ultrafine grain sizes cannot be achieved in pure Mg through the use of EX-ECAP: the minimum grain size in these experiments was ⬃26 m. By contrast, the presence of Zr particles inhibits grain growth in the Mg-Zr alloy, and it was possible to use EX-ECAP to achieve grain sizes of ⬃1.4 m. 4. Tensile testing at 523 and 573 K revealed enhanced ductilities in the Mg-Zr alloy, with maximum elongations up to ⬃300 to 400 pct. The measured values for the strainrate sensitivity under these conditions were ⬃0.3 to 0.4. This suggests that a dislocation-glide process may be ratecontrolling, but probably with some contribution from grain-boundary sliding because of the very fine grain size. ACKNOWLEDGMENTS We are grateful to Dr. Koichi Makii (Kobe Steel, Ltd., Kobe, Japan) for providing the material used in this investigation and we thank Takayoshi Fujinami for experimental assistance. This work was supported in part by the Light Metals Education Foundation of Japan, in part by the Mitsubishi Foundation, and in part by the United States Army Research Office under Grant No. DAAD19-00-1-0488. REFERENCES 1. F.H. Froes, D. Eliezer, and E. Aghion: JOM, 1998, vol. 50 (9), pp. 30-34. 2. K. Johnson: Adv. Mater. Proc., 2002, vol. 160 (6), pp. 62-65. 3. H. Friedrich and S. Schumann: J. Mater. Proc. Technol., 2001, vol. 117, pp. 276-81. 4. R.Z. Valiev, R.K. Islamgaliev, and I.V. Alexandrov: Progr. Mater. Sci., 2000, vol. 45, pp. 103-89. 5. V.M. Segal, V.I. Reznikov, A.E. Drobyshevskiy, and V.I. Kopylov: Russ. Metall., 1981, vol. 1, pp. 99-105. 6. V.M. Segal: Mater. Sci. Eng., 1995, vol. A197, pp. 157-64. VOLUME 35A, JUNE 2004—1743
10-03-80A-4.qxd
27/4/04
17:24
Page 1744
7. M. Furukawa, Z. Horita, M. Nemoto, and T.G. Langdon: J. Mater. Sci., 2001, vol. 36, pp. 2835-43. 8. N.A. Smirnova, V.I. Levit, V.I. Pilyugin, L.S. Davydova, and V.A. Sazonova: Fiz. Metall. Metalloved., 1986, vol. 61, pp. 1170-77. 9. A.P. Zhilyaev, S. Lee, G.V. Nurislamova, R.Z. Valiev, and T.G. Langdon: Scripta Mater., 2001, vol. 44, pp. 2753-58. 10. A.P. Zhilyaev, G.V. Nurislamova, B.-K. Kim, M.D. Baró, J.A. Szpunar, and T.G. Langdon: Acta Mater., 2003, vol. 51, pp. 753-65. 11. Y. Iwahashi, Z. Horita, M. Nemoto and T.G. Langdon: Acta Mater., 1997, vol. 45, pp. 4733-41. 12. Y. Iwahashi, Z. Horita, M. Nemoto, and T.G. Langdon: Acta Mater., 1998, vol. 46, pp. 3317-31. 13. J. Wang, Y. Iwahashi, Z. Horita, M. Furukawa, M. Nemoto, R.Z. Valiev, and T.G. Langdon: Acta Mater., 1996, vol. 44, pp. 2973-82. 14. Z. Horita. D.J. Smith, M. Furukawa, M. Nemoto, R.Z. Valiev, and T.G. Langdon: J. Mater. Res., 1996, vol. 11, pp. 1880-90. 15. P.B. Berbon, N.K. Tsenev, R.Z. Valiev, M. Furukawa, Z. Horita, M. Nemoto, and T.G. Langdon: Metall. Mater. Trans. A, 1998, vol. 29A, pp. 2237-43. 16. Y. Iwahashi, Z. Horita, M. Nemoto, and T.G. Langdon: Metall. Mater. Trans. A, 1998, vol. 29A, pp. 2503-10. 17. Z. Horita, T. Fujinami, and T.G. Langdon: Mater. Sci. Eng., 2001, vol. A318, pp. 34-41. 18. K. Nakashima, Z. Horita, M. Nemoto, and T.G. Langdon: Mater. Sci. Eng., 2000, vol. A281, pp. 82-87. 19. Y. Nishida, H. Arime, J.-C. Kim, and T. Ando: Scripta Mater., 2001, vol. 45, pp. 261-66. 20. A. Azushima and K. Aoki: Mater. Sci. Eng., 2002, vol. A337, pp. 45-49. 21. Y. Saito, H. Utsunomiya, H. Suzuki, and T. Sakai: Scripta Mater., 2000, vol. 42, pp. 1139-44. 22. J.-C. Lee, H.-K. Seok, J.-H. Han, and Y.-H. Chung: Mater. Res. Bull., 2001, vol. 36, pp. 997-1004. 23. J.-H. Han, H.-K. Seok, Y.-H. Chung, M.-C. Shin, and J.-C. Lee: Mater. Sci. Eng., 2002, vol. A323, pp. 342-47. 24. J.-C. Lee, H.-K. Seok, J.-Y. Suh, J.-H. Han, and Y.-H. Chung: Metall. Mater. Trans. A, 2002, vol. 33A, pp. 665-73. 25. J.-C. Lee, H.-K. Seok, and J.-Y. Suh: Acta Mater., 2002, vol. 50, pp. 4005-19. 26. M. Mabuchi, H. Iwasaki, K. Yanase, and K. Higashi: Scripta Mater., 1997, vol. 36, pp. 681-86. 27. M. Mabuchi, K. Ameyama, H. Iwasaki, and K. Higashi: Acta Mater., 1999, vol. 47, pp. 2047-57. 28. T. Mukai, M. Yamanoi, H. Watanabe, and K. Higashi: Scripta Mater., 2001, vol. 45, pp. 89-94. 29. H. Watanabe, T. Mukai, K. Ishikawa, and K. Higashi: Scripta Mater., 2002, vol. 46, pp. 851-56. 30. W.J. Kim, C.W. An, Y.S. Kim, and S.I. Hong: Scripta Mater., 2002, vol. 47, pp. 39-44.
1744—VOLUME 35A, JUNE 2004
31. A. Yamashita, Z. Horita, and T.G. Langdon: Mater. Sci. Eng., 2001, vol. A300, pp. 142-47. 32. Z. Horita, K. Matsubara, K. Makii, and T.G. Langdon: Scripta Mater., 2002, vol. 47, pp. 255-60. 33. K. Matusbara, Y. Miyahara, Z. Horita, and T.G. Langdon: Acta Mater., 2003, vol. 51, pp. 3073-84. 34. J.A. Chapman and D.V. Wilson: J. Inst. Met., 1962–63, vol. 91, pp. 39-40. 35. K.U. Kainer: Proc. 3rd Int. Magnesium Conf., G.W. Lorimer, ed., The Institute of Materials, London, 1997, pp. 533-43. 36. D.V. Wilson: J. Inst. Met., 1970, vol. 98, pp. 133-43. 37. M. Hilpert, A. Styczynski, J. Kiese, and L. Wagner: in Magnesium Alloys and Their Applications, B.L. Mordike and K.U. Kainer, eds., Wiley-VCH, Weinheim, Germany, 1998, pp. 319-24. 38. G.F. Weissmann and W. Babington: Proc. ASTM, 1958, vol. 58, pp. 869-86. 39. A.A. Nayeb-Hashemi and J.R. Clark: Binary Alloy Phase Diagrams, T. Massalski, ed., ASM, Metals Park, OH, 1986, pp. 1566-67. 40. Y. Iwahashi, J. Wang, Z. Horita, M. Nemoto, and T.G. Langdon: Scripta Mater., 1996, vol. 35, pp. 143-46. 41. C. Xu and T.G. Langdon: Scripta Mater., 2003, vol. 48, pp. 1-4. 42. M. Furukawa, Y. Iwahashi, Z. Horita, M. Nemoto, and T.G. Langdon: Mater. Sci. Eng., 1998, vol. A257, pp. 328-32. 43. M. Furukawa, Z. Horita, and T.G. Langdon: Mater. Sci. Eng., 2002, vol. A332, pp. 97-109. 44. T.G. Langdon: Metall. Trans. A, 1982, vol. 13A, pp. 689-701. 45. J. Wang, Z. Horita, M. Furukawa, M. Nemoto, N.K. Tsenev, R.Z. Valiev, Y. Ma, and T.G. Langdon: J. Mater. Res., 1993, vol. 8, pp. 2810-18. 46. Z. Horita, D.J. Smith, M. Nemoto, R.Z. Valiev, and T.G. Langdon: J. Mater. Res., 1998, vol. 13, pp. 446-50. 47. K. Oh-ishi, Z. Horita, D.J. Smith, and T.G. Langdon: J. Mater. Res., 2001, vol. 16, pp. 583-89. 48. S. Komura, Z. Horita, M. Furukawa, M. Nemoto, and T.G. Langdon: Metall. Mater. Trans. A, 2001, vol. 32A, pp. 707-16. 49. S. Lee, M. Furukawa, Z. Horita, and T.G. Langdon: Mater. Sci. Eng., 2003, vol. A342, pp. 294-301. 50. X. Wu and Y. Liu: Scripta Mater., 2002, vol. 46, pp. 269-74. 51. J.C. Tan and M.J. Tan: Scripta Mater., 2002, vol. 47, pp. 101-06. 52. W.-J. Kim, S.W. Chung, C.S. Chung, and D. Kum: Acta Mater., 2001, vol. 49, pp. 3337-45. 53. A. Wisbey and M.W. Kearns: in Superplasticity: 60 Years after Pearson, N. Ridley, ed., The Institute of Materials, London, 1995, pp. 305-23. 54. C.F. Dressel: in Superplasticity: 60 Years after Pearson, N. Ridley, ed., The Institute of Materials, London, 1995, pp. 359-76. 55. T.G. Langdon: Acta Metall. Mater., 1994, vol. 42, pp. 2437-43. 56. S.S. Vagarali and T.G. Langdon: Acta Metall., 1982, vol. 30, pp. 1157-70.
METALLURGICAL AND MATERIALS TRANSACTIONS A
