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Mg Alloy Rod Strengthened by Combined Processes of Deformation-Restricted Forging and Extrusion
2024 Volume 65 Issue 1 Pages 93-96
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2024 Volume 65 Issue 1 Pages 93-96
A new method for strengthening Mg alloy rod is proposed. AZ80Mg alloy rods were forged along the longitudinal axis (LA) at extremely high pressures beyond fracture stress under a condition where plastic deformation was eliminated in a die, i.e., deformation-restricted forging (DRFing), followed by cold extrusion. Although the ultimate tensile strength (UTS) and hardness were gradually raised with DRFing stress, the yield strength (YS) was lowered by sharp basal texture evolution on the plane normal to the LA and tensile axis. However, the extrusion after DRFing drastically changed the texture to (0001) || LA, causing a large increase in YS. Consequently, a superior balance of mechanical properties: YS of 376 MPa, UTS of 417 MPa, and ductility of 10%, could be achieved after the combined processes of DRFing and extrusion.
Stress-strain curves attained by tensile tests along the longitudinal direction of the processed AZ80Mg alloy rods; (a) as-DRFed rods and (b) DRFed rods followed by extrusion. For comparison, the flow curves of the as-hot extruded rod and it subjected to extrusion are also shown.
“Rod” is one of the most important forms of Mg alloys. They are machined to into products of any shapes. Nevertheless, research for strengthening Mg alloy rods is quite few. It should be due to the limited methods effectively applicable to rods for strengthening. Many methods have been attempted to strengthen Mg alloys in various forms; thus far, alloying by rare-earth addition1,2) and grain refinement by severe plastic deformation (SPD)3,4) are considered the most effective ones. Grain refinement is the classical and essential method for strengthening Mg alloys due to its relatively large Hall–Petch coefficients.5) However, given that the lower limitation of grain size $\bar{d}$ achievable by classical thermomechanical processes is $\bar{d}$ ≥ a few micrometers, SPD can promote further grain refinement to attain ultrafine-grained structures ($\bar{d} < 1$ µm).3,4,6,7) Nevertheless, the sample sizes fabricated by SPD methods are typically extremely small to be employed as structural elements.
Mechanical twinning is one of the most important mechanisms that causes grain refinement in Mg alloys.4,7,8) Different variants and types of twins necessary for grain fragmentation are frequently and densely developed under high deformation stress, exceeding their critical resolved shear stresses.4) However, applying excessive deformation stress to Mg alloys can cause failure, a paradox in fabricating bulky fine-grained Mg alloys by employing mechanical twinning. Hence, a limited number of reports carried out by SPD methods focusing on the ultrafine-grain evolution by mechanical twinning.4,9)
Recently, Miura et al. proposed a new method called “deformation-restricted forging (DRFing)”, in which Mg alloy rods were forged in a die under extremely high pressures beyond the fracture stress of rods.10) For a start, they demonstrated DRFing of AZ80Mg alloy “disks” under extremely high forging stresses over 900 MPa and superior balance of tensile properties; yield strength (YS) of 280 MPa, ultimate tensile strength (UTS) of 436 MPa, and a fracture strain of around 6% were derived by the tensile test along the direction normal to the forging direction. These mechanical properties were attributed mainly to the combined effects of large work hardening due to high dislocation density by extremely high applied forging stress, a sharp basal texture evolution parallel to the tensile axis (TA), i.e., (0001) || TA, and grain refinement by dense mechanical twinning. However, the evolved sharp texture conversely and seriously spoiled the YS along the forging direction. And they have not actually applied DRFing on to rods.
In the present study, hot-extruded AZ80Mg alloy “rods” were processed by DRFing and, then, they were followed by cold extrusion. The sharp basal texture normal to the longitudinal axis (LA) of rods, (0001) ⊥ LA, developed by DRFing should be destructed by the following extrusion to convert into another texture of (0001) || LA and, therefore, the YS along LA must be raised because of a mechanism of texture strengthening. Changes in the microstructure and mechanical properties during combined processes of DRFing and extrusion were precisely examined.
A commercial hot-extruded AZ80Mg alloy rod with a 19.6 mm diameter was cut into shorter pieces with a 50.0 mm length. Table 1 lists the chemical composition of the AZ80Mg alloy. The alloy pieces were placed in a cylindrical hole with a 20.0 mm diameter in a die followed by forging with applied flow stresses of 328 MPa to 984 MPa at an initial strain rate of 10−3 s−1 on an Amsler universal mechanical testing machine at room temperature (Fig. 1(a)). The die was made of SKD-11 steel water quenched and followed by tempering. When an initial AZ80Mg rod was deformed in compression without a die, fracture occurred at 409 MPa and strain of 12%. Thus, pressures considerably higher than the fracture stress in compression were applied to the rods during DRFing.10) The dimensions of the die hole were designed to achieve full contact of the rod with the die wall at a strain of approximately 4% in height, lower than the fracture strain. After full contact, the plastic deformation of the rod can proceed based on the elastic deformation of the die under a quasi-hydrostatic condition. As such, extremely high forging stresses can be successfully applied to the rod due to the restricted plastic deformation by the die. Although the additional amount of plastic strain after the full contact was limited and less than 1%, dislocation glide occurred to cause work hardening and further trigger mechanical twinning.10) After DRFing, the rods were extruded at room temperature, which showed around 6% reduction in diameter, i.e., about 12% plastic strain along the LA (Fig. 1(b)). The cumulative plastic strain along the LA after DRFing (4%) and extrusion (12%) was approximately 16%, larger than the compressive fracture strain.
Schematic illustrations of (a) DRFing of rods in a die and (b) cold extrusion. DRFing was followed by extrusion.
The evolved microstructure was investigated using orientation-imaging microscopy (OIM/TexSem Laboratories OIM ver. 5.0, Hitachi S-4500) at an accelerating voltage of 25 kV and transmission electron microscopy (TEM/Jeol 2000FX) at an accelerating voltage of 200 kV on the planes parallel to the LA of the rods after mechanical and electrical polishing. On the same planes, the change in hardness was also investigated using a micro-Vickers hardness tester. Tensile tests were carried out along the direction parallel to the LA using dog-bone-shaped specimens with 5.0 × 2.4 × 0.7 mm3 gauge dimensions on an Instron-type mechanical testing machine at an initial strain rate of 3.0 × 10−3 s−1 at ambient temperature.
Figure 2 shows the changes in the microstructure that evolved after DRFing. Grain fragmentation occurred by mechanical twinning. It is interesting to note in Fig. 1 that twins more densely evolved at high forging stresses. Hence, the average grain size of $\bar{d} = 10$ µm of the as-hot-extruded sample macroscopically decreased after forging at 984 MPa to approximately $\bar{d} = 5$ µm. Full contact of the rod with the die wall occurred at around 328 MPa. This result suggests that the dislocation motion necessary for mechanical twinning was still taking place in the rods under restrictions of plastic deformation during DRFing, probably due to elastic deformation of the die. More than 99% of the twins identified by OIM were {10–12} tensile types, and the others were rare. Because of the relatively low spatial resolution of the OIM system employed and the step size of 0.7 µm, substantially thinner compression twins are hardly detected. Miura et al. reported the TEM observations of the AZ80Mg alloy DRFed at 1003 MPa, indicating that nano twinning occurred in the coarse first-order mechanical twins.10) The actual grain size achieved after DRFing must be, hence, considerably finer than 5 µm. The intensity of the (0001) basal texture formed on the forging plane increased gradually with the increase in the forging stress from approximately 1.5 of the initial sample up to 6.7 after DRFing at 984 MPa. Therefore, a relatively sharp (0001) ⊥ LA texture developed on the forging plane of the rod by DRFing, suggesting that plastic deformation enabling crystallographical rotation still occurred under the constrained conditions in the die.
Microstructural changes before and after DRFing; (a) as-hot extruded samples and after forging at (b) 328 MPa, (c) 656 MPa, and (d) 984 MPa. An allow mark indicates the hot extrusion and the forging axes. (e), (f), (g), and (h) are the correspondent inverse pole figures of (a), (b), (c), and (d), respectively. Color decoding was changed along the forging axis for easy identification of the mechanical twins and change in the texture.
The DRFed rods were extruded along the LA at ambient temperature, and the evolved microstructures are displayed in Fig. 3. Almost all the evolved twins by DRFing (Fig. 2) disappeared after extrusion. Therefore, “detwinning” of the {10–12} tensile ones seemed to occur, and the fragmented grains by mechanical twinning appeared to be recovered as were those before DRFing. The developed (0001) ⊥ LA texture by DRFing (Fig. 2) completely disappeared. A sharp texture of (0001) || extrusion axis (EA) (i.e., equal to the LA) with intensity over 5, similar to that of the initial hot-extruded rod, was developed instead. Comparisons showed that the microstructure and texture almost exhibited the same independent process with or without DRFing in advance. Because of the dominant slip plane (0001) aligned parallel to the LA, slip deformation was difficult to occur during the tensile test along the LA. Therefore, texture strengthening as well as the effects of strain hardening was expected.10,11)
Microstructure and the texture evolved after the combined processes of DRFing and cold extrusion; Rods DRFed in advance at (b), (f) 328 MPa, (c), (g) 656 MPa, and (d), (h) 984 MPa. An allow mark indicates the cold extrusion axis, parallel to the LA and DRFing axis. (a) and (e) indicate the results after simple extrusion of the initial rod (i.e., hot extruded) specimen. Color decoding was changed parallel to the extrusion axis.
TEM observation was carried out on a sample processed by DRFing and extrusion, and a typical photograph is exhibited in Fig. 4. Figure 4(a) shows a fine acicular substructure with an appearance similar to fine twins4) with a boundary spacing of approximately 1 µm, which is considerably finer than the {10–12} twins detected by OIM after DRFing (Fig. 2). The selected-area diffraction pattern obtained from the center of Fig. 4(a) (indicated by the circle) is shown in Fig. 4(b). Analysis of the diffraction pattern revealed that the acicular substructure was composed of bands of the matrix and {10–12} twins (Fig. 4(c)). In addition, the matrix in acicular substructure exhibits a {10–12} twinning relationship with the region across the boundary, as indicated by the solid trace line in Fig. 4(a). Because the dotted and sold trace lines in Fig. 4(a) are symmetrical about the EA, the boundaries parallel to the dotted trace line must be also of {10–12} twins. Consequently, multiple twinning was confirmed. It is reported that multiple twinning largely contributes to homogeneous ultrafine-grained structure evolution in Mg alloys.4)
(a) Typical TEM micrograph of the sample forged at 984 MPa followed by extrusion. An arrow mark indicates the EA parallel to the LA. The dotted and solid lines are the traces of two different boundaries and the angles formed between EA and the traces. (b) Selected-area diffraction (SAD) pattern obtained from the area indicated by the circle in (a). (c) Analytical result of the SAD pattern of (b) indicating a {10–12} twinning relationship.
Detwinning occurs by reverse loading of Mg and Mg alloys.12,13) Morrow et al. demonstrated detwinning mechanism by in-situ TEM observation, showing that detwinning occurred through the re-nucleation and growth of second-generation twins. Therefore, the initial twins were replaced by the nucleation and growth of new ones at the sites of the previous ones. Miura et al. reported the TEM observation of DRFed AZ80Mg alloy, in which many nano twins were densely evolved in the coarse initial twins.10) In summary, the acicular substructure in Fig. 4 must be {10–12} tension twins formed by replacing initial twins. The microstructure obtained after DRFing followed by extrusion was therefore composed of acicular fine grains, which were extremely fine for identification by OIM.
Tensile tests of the processed samples were carried out and Fig. 5 shows the flow curves. The as-DRFed samples exhibited clear yielding at low-stress regions, followed by large strain hardening and large plastic deformation, indicating large plastic deformability in tension after DRFing (Fig. 5(a)). The clear yielding at the low-stress region and the large ductility in Fig. 5(a) were attributed to the sharp basal texture that developed normal to the TA, i.e., (0001) ⊥ TA. While the YS of the as-hot-extruded rods was 252 MPa, it after DRFing at 328 MPa was as low as 120 MPa despite a finer-grained structure than that of the former. However, the YS gradually increased to 172 MPa as the DRFing stress increased. The large plastic deformation region and low YS followed by appearance of high UTS were induced by the crystallographical rotation from (0001) ⊥ TA to (0001) || TA.14) Meanwhile, the flow curves of the samples processed by cold extrusion (Fig. 5(b)) appeared different from those of the as-DRFed samples in Fig. 5(a), indicating higher YS. The YS of the as-hot-extruded bar (252 MPa) was raised by cold extrusion up to 291 MPa. On the other hand, the YS of the rods processed by DRFing and extrusion further increased up to 376 MPa with the increase in applied stress during DRFing. However, the UTS was unchanged at roughly around 415 MPa. This should be lowered plastic elongation with increasing DRFing stress owing to grain refinement and larger work hardening during processing. Actually, hardness was gradually increased with DRFing stress and was further raised by extrusion up to approximately 900 MPa, while the corresponding figure are not shown here. The notable increase in the YS after extrusion with the increase in DRFing stress is due to grain refinement, strain hardening, and texture strengthening. Contrary to the increase in YS, the sharp (0001) || TA texture evolution hindered plastic deformability, and therefore, ductility decreased to approximately half of its original value. Yet, a superior ductility of around 10% was retained at the minimum. The results from Fig. 5 firmly indicate the important factors of texture strengthening for the mechanical properties of Mg alloys. The newly proposed combined processes of DRFing and extrusion successfully improved the mechanical properties of AZ80Mg alloy rods and balance of mechanical properties. In addition, YS of 376 MPa, UTS of 417 MPa, and ductility of 10% were achieved at best.
Stress - strain curves attained by tensile tests along the longitudinal direction of the processed AZ80Mg alloy rods; (a) as-DRFed rods and (b) DRFed rods followed by extrusion. For comparison, the flow curves of the as-hot extruded rod and it subjected to cold extrusion are also shown.
A few studies reported the improvement of the mechanical properties of AZ80Mg alloy by SPD. Miura et al. carried out multi-directional forging (MDFing) of AZ80Mg alloy at room temperature, producing acicular ultrafine-grained structures of about $\bar{d} = 0.3$ µm and excellent mechanical properties; YS of 530 MPa, UTS of 650 MPa, and 9% elongation to fracture were attained.4) They also conducted MDFing of AZ80Mg alloy under decreasing temperature conditions employing dynamic recrystallization, resulting in a homogeneous microstructural evolution with an average grain size of 0.8 µm and well-balanced mechanical properties, YS of 300 MPa, UTS of 445 MPa, and ductility of 10%–17%, could be achieved.7) Meanwhile, research on strengthening Mg alloys by adding rare-earth elements has been widely conducted. For example, Kawamura and Yamasaki reported notable strengthening by rare-earth addition; YS of around 380 MPa, UTS of approximately 400 MPa, and elongation of 5%.1) While the AZ80Mg alloy MDFed at room temperature exhibited outstanding mechanical properties, those of AZ80Mg alloy rods produced by the combined processes of DRFing and extrusion were on par with those attained by MDFing under decreasing temperature conditions and rare-earth addition.
A new process of deformation-restricted forging (DRFing) followed by cold extrusion, a facile method for forging Mg alloy rods in a die at a forging pressure largely exceeding the fracture stress, was proposed. The mechanical properties of AZ80Mg alloy rods can be successfully improved. By the combined processes of DRFing and extrusion, a superior balance of mechanical properties, yield strength of 376 MPa, UTS of 417 MPa, and ductility of 10% can be achieved. The yield strength after extrusion was distinctly raised with increasing DRFing stress. The developed coarse twins obtained by DRFing disappeared after extrusion, in which finer ones replaced the initial twins. The improvement of mechanical properties was derived by i) large strain hardening, which was enabled by controlling the basal texture during the DRFing and extrusion processes, ii) introducing dense fine mechanical twins due to extremely high forging stress, and iii) a sharp basal texture evolution parallel to the tensile direction.
The authors acknowledge the financial support given by the Light Metal Educational Foundation, Japan.
