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Seyed Moien Faregh, Ghader Faraji, Mahmoud Mosavi Mashhadi, and Mohammad Eftekhari, Texture evolution and mechanical anisotropy of an ultrafine/nano-grained pure copper tube processed via hydrostatic tube cyclic expansion extrusion, Int. J. Miner. Metall. Mater., 29(2022), No. 12, pp.2241-2251. https://dx.doi.org/10.1007/s12613-022-2514-4
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Owing to the malleability, ductility, and significant thermal and electrical conductivity of pure copper tubes, these parts have widespread use in industrial applications. On the other hand, it is well established that ultrafine grained (UFG) and nanostructured (NS) materials exhibit outstanding mechanical and physical properties [1]. Severe plastic deformation (SPD) methods warrant the attainability of said materials via grain refinement. Some materials produced by SPD methods possess an average grain size of 100 nm or less. Several popular SPD methods invented for the bulk materials are such as equal channel angular pressing (ECAP) [2–4], high-pressure torsion (HPT) [5], cyclic extrusion compression (CEC) [6], cyclic expansion extrusion (CEE) [7], and accumulative roll bonding (ARB) [8–9]. During the last two decades, much attention has been devoted to the development of SPD methods competent in production of UFG tubes. Some of these methods include tubular channel angular pressing (TCAP) [10], parallel tubular channel angular pressing (PTCAP) [11–12], tube channel pressing (TCP) [13], cyclic flaring and sinking (CFS) [14], and high-pressure tube twisting (HPTT) [15]. Almost all of the methods designed during these years suffer from one critical flaw: their ineptitude in fabricating long and large-scale tubes. The main reason for this is the increasing amount of friction force. By increasing the length of the workpiece, the contact surface between the tube and die increases, and as the friction force rises the pressing punch starts to buckle under high amounts of total force. In 2018, Motallebi Savarabadi et al. [16–17] introduced an innovative SPD method entitled hydrostatic tube cyclic expansion extrusion (HTCEE). In this method, owing to existence of hydraulic fluid, the friction force is independent from tube’s length. Hence, making the production of a long tube possible. A schematic of the die setup is shown in Fig. 1(a). The process consists of two half-cycles. During the first one, the tube and the moving mandrel inside it are rammed into the die chamber using a pressing punch and because of the existence of back-pressure system, the tube expands in the deformation zone (Fig. 1(b)). The fluid filling the space between the tube and die is sealed using a poly tetra fluoro ethylene (PTFE) polymeric seal. In the next half-cycle, the back-pressure system is removed, and with the help of the moving mandrel and the PTFE seal the tube is extruded back to its initial shape, marking the end of the first cycle (Fig. 1(c)). In the following cycle, the whole die is rotated and each half-cycle is repeated on the other side of the tube (Fig. 1(d)). Motallebi Savarabadi et al. [17] saw that after applying two passes of HTCEE on the commercially pure copper tube, more refined and homogeneous microstructure was obtained compared to the as-received and the one-pass HTCEE processed samples. Also, they observed that the yield strength (YS) increased from 75 to 310 MPa, and the ultimate tensile strength (UTS) and microhardness were enhanced from 207 to 386 MPa and from HV ~59 to HV ~143, respectively. But during this process, elongation to failure decreased from ~55% to 37%. In other words, a notable increase in the hardness and strength was achieved besides a low loss of ductility. After N passes of the HTCEE process the equivalent plastic strain (
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εp=2N[ln(R2d−r2R2−r2)+4√3cot(α2)] |
(1) |
The substantial amount of plastic strain induced during SPD methods also leads to texture development. Evolution of texture and the grain refinement occurring during plastic deformation could influence the anisotropic behavior of the polycrystalline metal. Various researches have been conducted on the plastic and mechanical anisotropy of nanostructured materials. However, only a handful of them aim at analyzing this phenomenon in tubular components. As a case in point, Tavakkoli et al. [18] studied the mechanical anisotropy of nanostructured Cu–Zn tubes processed by PTCAP method. They reported that there was a severe anisotropy in UFG tubes in comparison to their course grained counterparts. In their study, a high strength of ~1220 MPa along the peripheral direction was observed while this value along the axial direction was ~580 MPa. Ultimate strength was enhanced ~106% and 66% along the peripheral and axial directions, respectively. Also, they reported that after processing by PTCAP, the {111} plane was rotated ~90ο with reorientation toward the radial direction, and this event is the main reason of the observed lower yield stress in the axial direction in comparison with the peripheral direction. Faraji et al. [19] made use of the concept of R-value (as an indicator of plastic anisotropy) to describe the state of directionality in an UFG Al6061 tube’s properties. It was realized that R-value relies heavily on the amount of equivalent plastic strain and also, as the texture weakens due to grain refinement, it decreases. Concerning non-tubular materials, Zhang et al. [20] investigated the anisotropy of a ZK61 sheet fabricated by extrusion and multi–pass rolling. They observed that grain refinement has an alleviation effect on mechanical anisotropy. In this way, after performing hot rolling, planar texture anisotropy and microstructure homogeneity was considerably improved due to multiple dynamic recrystallization. A uniform fine-grained microstructure with an average grain size of 6.1 µm and a {0002} basal texture was attained leading to enhanced strength isotropy in the plane and also high YS, which is attributed to pyramidal (<c+a>) slip. Sabirov et al. [21] studied the anisotropic behavior of a bulk nanostructured pure Ti produced by a combination of metal forming methods. The enhancement of mechanical behavior and reduction of mechanical anisotropy in a magnesium alloy processed by equal channel angular processing (ECAP) was studied by Suh et al. [22]. Texture evolution and anisotropy in the thermo-mechanical response of UFG Ti processed via ECAP was investigated by Meredith and Khan [23]. In their study, the sheet was ECAPed along the rolling direction, then it was ECAPed again but this time along the transverse direction by rotating 90° in the sheet plane. This procedure resulted in grain refinement and developed basal texture with tilted basal planes towards the direction of pressing. Also, the presented procedure could reduce mechanical anisotropy and enhance hardening behavior. Finally, an almost isotropic hardening at room temperature with improved ductility was reported for AZ31 sheet processed by ECAP through the mentioned procedure. Al-Maharbi et al. [24] presented the idea that with proper selection of starting texture it is possible to control the evolution of plastic anisotropy and reach intensely or weakly anisotropic mechanical properties in Mg alloys during equal channel angular extrusion (ECAE) processing. Sometimes, researchers use the metal forming methods to produce UFG metal with improved mechanical properties and special texture condition. As a case in point, Mao et al. [25] utilized rotary swaging process for producing UFG copper rods. They reported that after the swaging process by equivalent plastic strain of 2.5, the initial 54 μm coarse grains were gradually elongated into long columnar grains with a mean diameter of 2.06 μm and a length of 339 μm. Also, they saw that the swaging process led to strong <111> and weak <100> fiber textures along the Cu axis. Also, this process resulted in an increase in the YS from 60 to 450 MPa and a decreases in the ductility from 57% to down to 10% when ε = 2.5. In other work, Meng et al. [26] studied the microstructure and mechanical properties of commercial pure titanium subjected to rotary swaging. They could produce UFG titanium rod with superior mechanical properties. Also, Wan et al. [27] and Chen et al. [28] produced bulk nanostructured magnesium alloys using rotary swaging process.
In this research, UFG pure copper tubes were produced using the HTCEE process. Then, the effects of HTCEE process on the anisotropic behavior, texture, and mechanical properties of the coarse-grained and UFG tubes were characterized and compared along longitudinal and circumferential directions. In this way, tensile testing in two directions, texture analysis by X-ray diffraction (XRD) and evaluation of {111} pole figures, optical microscopy (OM) evaluations, hardness measurements, and fractography analysis were done.
Commercially pure copper (Cu ~99.90%) tubes were machined to an outer diameter, thickness, and length of 30 mm, 3.5 mm, and 100 mm, respectively. Before HTCEE process, the tubes were annealed at the temperature of 600°C for a time period of 2 h [29] leading to an almost homogeneous recrystallized microstructure. A HTCEE die with the following geometric parameters (as shown in Fig. 1(b)) was used: r = 11.5 mm, R = 15 mm, Rd = 17 mm, Rfillet = 3 mm, and α = 150°. The PTFE seal was designed marginally larger than the die’s inner diameter in order to prevent leakage during processing. The HTCEE process applies an equivalent plastic strain of 1.66 in each cycle (calculated by Eq. (1)). The deformation process was performed at 5 mm·min−1 ram speed at room temperature. Fig. 2 demonstrates the as-annealed and HTCEE processed tubes. OM was used to characterize the microstructure of the metal in as-annealed condition and in both orthogonal directions after HTCEE processing. The samples were mounted and polished before imaging. To scrutinize the microstructure evolution that occurred after performing HTCEE process, transmission electron microscopy (TEM) was used. The TEM specimen were prepared from the cross section perpendicular to the tube axis using the in situ lift-out procedure on a dual beam FEI Nova Nanolab 600 system. Longitudinal and peripheral tensile test samples having a gauge length of 9.6 and 24 mm, respectively, were prepared using wire electro discharge machining and tested in tension at ambient temperature and a strain rate of 10−3 s−1. Fig. 3 demonstrates the schematic of the ring hoop tensile test fixture and the position of elements employed for microstructure, texture, and hardness measurements. Multiple layers of oil were used in order to reduce the friction between split disks and the peripheral specimen. The Vickers hardness measurements were recorded in both longitudinal and peripheral directions at room temperature under a load of 0.49 N with a dwell time of 15 s for each separate measurement. The texture of the as-annealed and HTCEE processed tubes were analyzed using XRD and {111} pole figures were obtained on both RD/ED and RD/TD planes where RD, ED, and TD were defined in Fig. 3. After tensile testing, to study the mechanisms of fracture of specimens, fracture surfaces were analyzed using scanning electron microscopy (SEM).
Fig. 4 shows OM images of the pure copper in the as-annealed condition and after HTCEE processing through one cycle in longitudinal direction and peripheral direction. According to Fig. 4(a), a nearly equiaxed microstructure is obseved in the as-annealed sample, but after a cycle of HTCEE (
Stress–strain curves of the processed and as-annealed pure copper tubes in orthogonal directions are shown in Fig. 6(a). Based on the well-known Hall–Petch relationship, see Eq. (2), the mechanical properties and deformation behavior of the materials depends on grain size [36]. The values of UTS and YS increase in both directions due to the grain refinement occurring as a result of severe plastic deformation while the values of elongation to failure (El) undergo a slight reduction [37]. Similar behavior of the enhancement of material strength, after applying SPD methods, is also seen in other researches [2,11–12,17,30–31]. The main reasons for the enhancement of the yield and ultimate strength are the dislocations strengthening and the grain boundaries strengthening. In dislocation strengthening (or strain hardening), a high density of dislocations is produced which leads to the enhancement of the strength in the severely deformed material. Grain boundaries strengthening is another reason for the strength enhancement. In this condition, by applying severe plastic strain, in the presence of high hydrostatic pressure which commonly exists in SPD processes, an intense decrease in the grains sizes occurs (as argued in previous section). In the next stage, the grain boundaries and twins play the role of an obstacle on the path of the dislocations movement. This event leads to an enhancement in the resistance to deformation of the material and consequently, leads to the further enhancement of strength [18]. In this way, subsequent passes of SPD processes can lead to further enhancement of the strength [2,11,17].
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σ0=σi+kD−12 |
(2) |
where
According to Fig. 6(b), the UTS increases to 345 and 401 MPa along axial and peripheral directions, respectively. After HTCEE processing, the YS was enhanced to 270 MPa and 325 MPa along axial and peripheral directions, respectively. As it is obvious, the values of UTS and YS in the axial direction is lower than those in the peripheral direction, and it is in agreement with the results of Fig. 4 because, as seen in Fig. 4(b) and (c), the HTCEE processed material experiences more grain refinement in the peripheral direction than the axial direction. Thus, according to Hall–Petch relationship presented in Eq. (2), higher mechanical properties are expected in the peripheral direction. Another reason could be explained by the fact that, as a face centered cubic (FCC) metal undergoes deformation, some slip bands are formed that will orient in specific directions. Meaning that their activation occurs only if the force is applied at that particular angle [38]. When these slip bands are activated in a particular direction, dislocations movement is achieved more effortlessly and therefore the UTS and YS of the material will be smaller in some directions.
As was demonstrated in Fig. 6, different deformation behavior and tensile properties were recorded for the specimens along axial and peripheral directions. The ratio of UTS in the peripheral direction to that in the axial direction could be defined as a criterion for mechanical anisotropy. These values are 1.7 and 1.16 for the as-annealed coarse-grained and HTCEE processed ultrafine-grained tube, respectively. It is obvious that HTCEE process has had a reducing effect on the mechanical anisotropy of the pure copper tube. This phenomenon could be attributed to the relatively high SFE of pure copper (78 mJ/m2) and the mechanism of migration and multiplication of dislocations which gets activated during plastic deformation. Javidikia and Hashemi [39] observed the same trend in an aluminum alloy severely deformed via PTCAP. Defining the same criteria for mechanical anisotropy, it was reported that severe plastic deformation had reduced the mechanical anisotropy from a value of 2 to 1.73. The high SFE of aluminum alloy 5083 and the activation of slip mechanisms were cited as a reason. On the other hand, Tavakkoli et al. [18] observed that plastic deformation has had a worsening effect on the mechanical anisotropy of brass tubes, increasing it from an initial value of 1.69 to 2.1. The fact that brass is considered a low SFE material and deforms mostly by the twinning mechanism were held accountable. In materials having low SFE, the mechanism of twinning is preferred to the slip because of the lower activation energy of twinning. The distance of free slip line decreases due to the twin distributions. In Table 1, the value of mechanical anisotropy obtained from this study is compared with the results of other studies [18,39–40]. From Table 1, the HTCEE process leads to a significant reduction in the mechanical anisotropy, which is one of the main advantages of HTCEE process.
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In addition to severe tensile and YS anisotropy, the ductility of the specimens is also quite different along orthogonal directions. By comparison to as-annealed tube, the elongation to failure of HTCEEed tube decreases about 14% along the axial direction and 9.7% along the peripheral direction. Simply put, there is less formability loss in peripheral direction in comparison to axial direction. This may be as a result of the particular orientation of the slip systems during the tensile tests, which influences the ductility of the material [41]. In Table 2, the UST and El to failure, along the axial direction, and also hardness of 1 pass HTCEE processed tube are compared with those of other studies [29,42–50] performed on pure copper. From Table 2, in comparison to other techniques, 1 pass HTCEE shows a higher value of El beside high values of strength and hardness. This feature is another important advantage of HTCEE process. It was reported that HTCEE process contains higher levels of hydrostatic compressive stresses [17]. These stresses beside the accumulated plastic strain and shear strain, play an important role in the grain refinement occurred in SPD processes. The higher hydrostatic compressive stresses, by delaying the crack initiation, closing the cracks, and also limiting the cracks growth, can lead to the improvement of the workability of the material to achieve higher strains before early failure caused by the cracks initiation and propagation. This leads to the prevention of significant loss of ductility [51–52]. Thus, HTCEE process can result in the higher strengths and the lower losses of ductility [17].
| Number | SPD process | UTS / MPa | Hardness, HV | El / % | Ref. |
| 1 | 1 pass HTCEE | 345 | 132 | 40 | This study |
| 2 | 1 pass ECFE | 245 | 109 | 15 | [41] |
| 3 | 1 pass ARB | 357 | 112 | 2.5 | [42] |
| 4 | 1 pass ARB | 390 | 115 | ~5 | [43] |
| 5 | 1 pass ECAP | 314 | — | 9.5 | [44] |
| 6 | 1 pass RF | 295 | 70 | 18 | [45] |
| 7 | 1 pass twist CAP | 318 | 105 | 28 | [46] |
| 8 | 1 pass PTCAP | 430 | — | 37 | [47] |
| 9 | 1 pass HPTE | — | 121 | — | [48] |
| 10 | 1 pass TCEC | 273 | 80 | 18 | [49] |
| 11 | 1 pass ECAP | — | 115 | — | [50] |
| Notes: ECFE—equal channel forward extrusion; RF—repetitive forging; CAP—twist channel angular pressing; HPTE—high pressure torsion extrusion; TCEC—tube cyclic extrusion compression. | |||||
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The values of the Vickers microhardness, HV, are plotted in Fig. 7 against the distance from the inner surface of the as-annealed and UFG tubes in both directions. It is apparent from Fig. 7 that hardness has a remarkable increase after a single cycle of HTCEE. The microhardness of the as-annealed sample, which is HV 59, is increased to about HV 132 after one cycle of HTCEE processing. The escalation observed in hardness values is directly associated with the decrease in grain size and the increase of dislocations densities [45]. Concerning Fig. 4, the HTCEE processed tube, which experiences more grain refinement in both directions, possesses higher values of hardness compared to the as-annealed sample. This is in agreement with Hall–Petch relationship for hardness [53], which denotes that the values of hardness and the grain size of the material have an inverse relationship. A similar behavior of the enhancement of microhardness after SPD methods is also reported in other studies [2,11–12,17]. This trend is caused by the grain refinement, an enhancement of the grain boundaries amount, an increase in the dislocations density, and work hardening [2,51,54]. It is reported that by applying further passes of SPD process, microhardness and dislocations density are finally saturated. In this condition, the steady-state density of dislocations exist, which is as a result of establishing a balance between dislocation creation during SPD process and annihilation due to dynamic restoration phenomena [43]. As depicted in the diagram of Fig. 7, hardness in circumferential direction reaches slightly higher values than that in axial direction. The results from tensile tests demonstrated an identical trend in strength values. As mentioned earlier, in Table 2, the average value of hardness of 1 pass HTCEE processed tube are compared with that of other studies performed on pure copper. It was observed that 1 pass HTCEE leads to a significantly high value of hardness.
Material texture is a microstructural property which describes the distribution of crystallographic orientations of a polycrystalline sample. SPD processes can lead to significant changes of texture. It can be related to applying higher shear strain to the material and the direction of material flow during process, as mostly seen in processes such as ECAP and HPT. The post-deformation texture can influence and determine many characteristics of the material’s behavior including plastic anisotropy, strength, formability, fracture mechanism, and grain refinement [55]. Grain refinement by severe plastic deformation is frequently associated with texture evolution. As a polycrystalline metal undergoes deformation, the crystal lattices of the newly formed grains reorient in a preferred direction and lead to the evolution of texture. For texture analysis, {111} pole figures of the as-annealed and HTCEE processed samples were obtained in both directions using XRD technique; the corresponding results are presented in Fig. 8. As was anticipated, plastic deformation has a vigorous effect on texture formation and the post-deformation textures (Fig. 8(b) and (d)) appear to be completely different from the recrystallized ones (Fig. 8(a) and (c)). {111} pole figure in the axial direction of the as-annealed sample (Fig. 8(a)) demonstrates a relatively low texture intensity and a random texture along RD. After HTCEE processing, the oriental distribution of grains barely changes however a fair amount of increase in the texture intensity is observed (Fig. 8(b)). The reason for this phenomenon could be the effects of new grains generated by grain refinement and continuous grain recrystallization due to the accumulation of strain, and also the grain rotation during the SPD process [55–56]. Fig. 8(c) and (d) depicts the {111} pole figures of the as-annealed and HTCEE processed specimens in peripheral direction. In the as-annealed condition, a random texture along TD is observed but after plastic deformation the poles reorient mostly along the RD, explaining the higher values of YS in peripheral direction. Both as-annealed and HTCEE processed sample show higher texture intensity in peripheral direction than axial direction. Similar trend of the observed texture intensity was also seen in the {111} pole figures of Ref. [18], which is related to the mechanical anisotropy of nanostructured Cu–Zn tubes processed by PTCAP process.
Fig. 9 demonstrates the SEM micrographs of the fractured surfaces after tensile testing. The size, shape, and orientation of the dimples determines the mode of fracture. As can be seen in Fig. 9(a), a large number of deeper, coarser, and more equiaxed dimples signifying ductile fracture mode exist on the fracture surface of the as-annealed specimen compared to the processed sample, and these features is indicative of more formability of this sample. The dimples size reduction after applying SPD methods is a common happen reported in other studies [2,17,29]. This is attributed to the grain refinement and work hardening happened during SPD methods. In other words, as is obvious in Fig. 6(a), after applying HTCEE procedure, the capability of work hardening is reduced. This leads to the reduction of the values of uniform elongation and ductility. Thus, after the nucleation stage, the dimples do not have adequate time to grow and coalescence with each other, which is resulted in the shallower dimples [57]. It was also reported that further passes of SPD methods can lead to the smaller and shallower dimples [2,17]. Fig. 9(b) demonstrates the SEM fractograph of the processed sample in longitudinal direction. As depicted, after HTCEE processing a decrease in the diameter and depth of the co-axial micro voids or dimples oriented along the rolling direction is observed. These shallow dimples accompanied by a higher amount of tearing edges indicate the occurrence of ductile fracture mode [58–59], and lead to the reduction of formability. This kind of fracture happens by microvoid formation, coalescence with each other and consequently, crack formation and then, the crack propagation [58]. Moreover, some indications of cleavage planes and brittle fracture can be recognized in the HTCEE processed specimen. Further passes of HTCEE process may change the fracture mechanism from ductile to brittle. SEM fractograph of the processed sample in the peripheral direction is shown in Fig. 9(c). This figure as well illustrates that the dimples undergo a decrease in diameter and depth after processing. However, in this direction, the dimples orient along the thickness direction. It seems that after HTCEE, the specimen fractures by axial tension mode inlongitudinal direction and shear mode in peripheral direction.
Texture evolution and mechanical anisotropy of an ultrafine/nano-grained pure copper tube processed by hydrostatic tube cyclic expansion extrusion was studied through experimental procedures, and several astonishing conclusions were made.
(1) HTCEE process via the grain refinement and the formation of an UFG microstructure (grain size mostly less than 200 nm) leads to the enhancement of mechanical properties including, strength and hardness in both axial and peripheral directions.
(2) Based on tensile tests results, a different deformation behavior is observed along axial and peripheral directions. The UTS was enhanced to 345 and 401 MPa along axial and peripheral directions, respectively. The YS increased to 270 and 325 MPa along axial and peripheral directions, respectively. Dislocation strengthening (or strain hardening) and grain boundary strengthening are two main reasons for the enhancement of strength in HTCEE processing of pure copper.
(3) El to failure of the UFG specimen decreases about 14% along the axial direction and 9.7% along the peripheral direction in comparison to the as-annealed specimen. After HTCEE process, a low loss of ductility was observed in both directions, which is one of the advantages of HTCEE process. This is attributed to the higher hydrostatic compressive stresses of the HTCEE process, which by delaying the crack initiation, closing the cracks and also limiting the cracks growth, can lead to the improvement of workability of the material to achieve higher strains before early failure. This leads to the prevention of significant loss of ductility.
(4) The ratio of UTS in the peripheral direction to that in the axial direction, as a criterion for mechanical anisotropy, are 1.7 and 1.16 for the as-annealed coarse-grained and HTCEE processed UFG tube, respectively. This indicates that the HTCEE process led to a reduction in the mechanical anisotropy of the pure copper tube, which is another advantage of HTCEE process.
(5) Hardness has a remarkable increase after a single cycle of HTCEE (from HV 59 to HV ~132). This behavior is caused by the grain refinement, an enhancement of the grain boundaries amount, an increase in the dislocations density, and work hardening. Also, similar to the tensile strength of the metal, the Vickers microhardness value in peripheral direction is greater than that in longitudinal direction.
(6) XRD measurements confirm the anisotropy observed in the mechanical behavior of the material.
(7) Different modes of fracture were observed after tensile testing in both orthogonal directions. But, according to the fracture evidences, it seems that ductile fracture occurred predominately in all tensile specimens. Also, after HTCEE, the dimples undergo a decrease in diameter and depth. This is attributed to grain refinement and work hardening happened during SPD methods. In this condition, the dimples do not have adequate time to grow and coalescence with each other.
The authors declare no conflict of interest.
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| Number | SPD process | UTS / MPa | Hardness, HV | El / % | Ref. |
| 1 | 1 pass HTCEE | 345 | 132 | 40 | This study |
| 2 | 1 pass ECFE | 245 | 109 | 15 | [41] |
| 3 | 1 pass ARB | 357 | 112 | 2.5 | [42] |
| 4 | 1 pass ARB | 390 | 115 | ~5 | [43] |
| 5 | 1 pass ECAP | 314 | — | 9.5 | [44] |
| 6 | 1 pass RF | 295 | 70 | 18 | [45] |
| 7 | 1 pass twist CAP | 318 | 105 | 28 | [46] |
| 8 | 1 pass PTCAP | 430 | — | 37 | [47] |
| 9 | 1 pass HPTE | — | 121 | — | [48] |
| 10 | 1 pass TCEC | 273 | 80 | 18 | [49] |
| 11 | 1 pass ECAP | — | 115 | — | [50] |
| Notes: ECFE—equal channel forward extrusion; RF—repetitive forging; CAP—twist channel angular pressing; HPTE—high pressure torsion extrusion; TCEC—tube cyclic extrusion compression. | |||||
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