Journal of Materials Research and Technology Journal of Materials Research and Technology
Original Article
Direct influence of recovery behaviour on mechanical properties in oxygen-free copper processed using different SPD techniques: HPT and ECAP
Meshal Y. Alawadhia,, , Shima Sabbaghianradb, Yi Huanga,, , Terence. G. Langdona
a Materials Research Group, Faculty of Engineering and the Environment, University of Southampton, Southampton SO17 1BJ, UK
b Department of Chemical Engineering and Materials Science, University of Southern California, Los Angeles, CA 90089-0241, USA
Received 11 April 2017, Accepted 15 May 2017
Abstract

Oxygen-free copper of 99.95wt.% purity was severely deformed at room temperature by two modes of severe plastic deformation, equal-channel angular pressing (ECAP) and high-pressure torsion (HPT). ECAP was performed using 4, 16 and 24 passes, and HPT was performed using 1/2, 1 and 10 turns. The results show that while recovery occurs during both ECAP and HPT processing, copper shows a faster recovery rate with HPT processing than ECAP. The occurrence of recovery was observed at an equivalent strain exceeding ∼12 that led to an enhancement in the uniform plastic deformation. The influence of recovery behaviour on the mechanical properties was investigated using X-ray diffraction, microhardness and tensile testing.

Keywords
Ductility, Equal-channel angular pressing, High-pressure torsion, Recovery, Strain rate sensitivity, Work hardening
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1Introduction

High-pressure torsion (HPT) [1–4] and equal-channel angular pressing (ECAP) [2,5–7] are the most commonly used severe plastic deformation (SPD) techniques. The HPT procedure is widely used because it has the ability to produce exceptional grain refinement [4,8] and a large fraction of high-angle grain boundaries [9], while ECAP is popular because it is a simple process that has the ability to deform larger samples which makes it a better candidate for structural applications [10–12].

It is well established that bulk ultrafine-grained (UFG) materials are successfully produced by SPD with very small grain sizes in the range between 1μm and 100nm [2,13–16]. UFG materials normally exhibit an extraordinary increase in strength in comparison to their coarse grain counterparts [17,18]. The inverse relationship between the yield stress (σy) and the square root of the grain size (d) is described by the Hall–Petch relationship [19,20]:

where σ0 is the lattice friction stress and ky is a yield constant. It is readily apparent from Eq. (1) that the strength of the material increases as a consequence of the reduction in grain size. Several studies have reported the achievement of a significant grain refinement in pure Cu using ECAP [21–29] and HPT [30–41]. To date, most studies have suggested that pure Cu shows hardening behaviour without recovery during both HPT [31,35,41] and ECAP [42,43] at room temperature. Nevertheless, two recent studies have reported a softening behaviour of pure Cu with recovery during HPT processing [44,45]. This softening is a well-documented behaviour of pure Al [46–54] and has been observed in other materials such as pure Zn [53] and pure Mg [55].

A third recent study also reported the occurrence of a recovery mechanism in pure Cu that contributed to the enhancement of ductility during 1–16 passes of ECAP at room temperature [25]. This latter report confirmed the results published in the classic study [43] that high ductility was achieved after 16 passes of ECAP as a result of grain boundary sliding. These findings were unexpected because it was previously claimed that the increase in strength in UFG materials is associated with a decrease in ductility at ambient temperature [18,56], and the poor ductility of UFG materials is due to the limited strain hardening ability [18,57,58].

Two different deformation mechanisms are proposed for the simultaneous gain in strength and ductility after processing pure Cu by ECAP for 16 passes at room temperature: one is grain boundary sliding [43] and the other is a recovery mechanism [25]. The present study is designed to clarify the mechanism leading to the enhancement in ductility by processing oxygen-free Cu for 24 passes at room temperature. Twenty-four passes was chosen instead of 16 to impose similar maximum strain on the specimen because the die angle used in this study was 110°, whereas 90° was used in the previous studies. In addition, it was possible to investigate the potential for achieving high strength and ductility by imposing very high strain on the specimens using HPT since imposing a higher strain by ECAP was limited because it was difficult to process the specimens beyond 24 passes due to the initiation of cracks.

2Experimental materials and procedures

Oxygen-free Cu (99.95%) was used in this investigation. Billets of length 65mm and 10mm diameter were processed by ECAP at room temperature. These billets were pressed through a solid die having an internal channel angle, Φ, of 110° and an outer arc of curvature, Ψ, of 20°. An equivalent strain of ∼0.76 was created on each separate pass based on the Φ and Ψ values [59]. The billets were pressed for 4, 16 and 24 passes repetitively and provided a maximum equivalent strain of ∼18 using route BC. In this route, the samples are rotated by 90° in the same direction between each separate pass [60]. It is well documented that route BC is the best choice for ECAP processing to give an array of UFG microstructure with equiaxed grains separated by a high fraction of high-angle boundaries [61].

HPT processing was conducted at room temperature using disc samples with 10mm diameter and ∼0.83mm thickness. The discs were compressed between two anvils under an applied pressure of 6.0GPa and a torsional strain was imposed by rotating the lower anvil at a speed of 1rpm. The discs were processed through 1/2, 1 and 10 turns using quasi-constrained conditions [62,63].

Prior to ECAP and HPT processing, the samples were annealed within a vacuum tube furnace at 600°C for 1h. The average grain size in the annealed sample was ∼24μm and the average Vickers microhardness was ∼41Hv.

The grain structure of oxygen-free Cu was examined by electron backscatter diffraction (EBSD) using a JEOL JSM-7001 F analytical field emission scanning electron microscope. An operating voltage of 15kV was used during the scanning process and a step size of 0.05μm was used to collect the EBSD patterns. The OIM images for HPT discs were taken at a distance between ∼3.0mm and ∼4.0mm from the disc centre.

The X-ray diffraction (XRD) testing was conducted on the samples after ECAP and HPT processing. A Bruker D2 Phaser X-ray diffractometer was used to analyze the whole surfaces of the samples using a copper target with Cu Kα (λ=0.15406nm) radiation. Recording the XRD patterns was performed by θ–2θ scans from 2θ=40° to 100° and profile fitting was accomplished by Maud software. Crystallite size and microstrain were calculated based on the Rietveld method using Maud software and their values were used to calculate the dislocation density.

Vickers microhardness measurements were recorded using a Future-Tech FM-300 microhardness tester. 100gf was used during the hardness indentations with a 15s dwell time. For HPT, the average values were recorded along the radius (5.0mm) of each disc with 0.3mm between each indentation point. Four indentations were measured around each of these points, separated by a distance of 0.15mm, then the average for these four points was calculated. For ECAP, the microhardness measurements were recorded along the longitudinal axis of the ECAP billets over a distance of 25mm with 0.5mm separation between each indentation point.

For tensile testing, tensile specimens were machined from HPT discs and ECAP billets using electrical discharge machining. The HPT specimen had a gauge length of 1.0mm and cross-sectional area of 1.0mm×0.8mm as shown in Fig. 1(a) and the ECAP specimen had a gauge length of 4.0mm and cross-sectional area of 3.0mm×2.0mm as shown in Fig. 1(b). A Zwick Z030 testing machine was used for pulling the specimens at room temperature using initial strain rates of 1.0×10−4s−1 and 1.0×10−3s−1.

Fig. 1.
(0.26MB).

Schematic illustration showing the dimensions of the miniature tensile specimens cut from (a) ECAP billet and (b) HPT disc.

3Experimental results3.1Microstructure after deformation

Fig. 2 presents OIM images for oxygen-free copper after processing by ECAP for 4 and 24 passes and HPT for 1/2 and 10 turns. The difference in colours represents the difference in the grain misorientations as denoted by the unit triangle.

Fig. 2.
(1.99MB).

EBSD orientation images showing the evolution of microstructure in oxygen-free copper specimens after ECAP processing through (a) 4 passes, (b) 24 passes and after HPT processing through (c) 1/2 turn and (d) 10 turns.

After 4 passes of ECAP, the grains were mostly large and elongated with a scattering of small grains as shown in Fig. 2(a). The average grain size was ∼4.5μm and the fraction of high-angle grain boundaries (HAGBs) was ∼65%. HAGBs are defined as boundaries with misorientations larger than 15°, whereas boundaries with misorientations between 2° and 15° are defined as low-angle grain boundaries (LAGBs). A significant grain refinement was observed after 24 passes as displayed in Fig. 2(b). The microstructure evolved into reasonable homogeneity with ultrafine and equiaxed grains having an average grain size of ∼600nm and ∼90% of HAGBs.

After 1/2 turn of HPT shown in Fig. 2(c), a rapid evolution towards microstructural homogeneity was observed. The microstructure consisted of ultrafine and equiaxed grains with an average size of ∼700nm and ∼85% of HAGBs. A further grain refinement was observed after deforming the disc with 10 turns as displayed in Fig. 2(d). High strain deformation produced an average grain size of ∼500nm with ∼82% of HAGBs.

3.2XRD analysis

The values of the calculated dislocation densities and crystallite sizes obtained from XRD measurements are shown in Fig. 3. The dislocation density was calculated based on the data obtained from XRD analysis using the equation [64,65]:

where ε21/2 is the lattice microstrain, DC is the average crystallite size and b is the Burgers vector. The XRD analysis on the annealed condition showed a dislocation density of ∼3.9×1012m−2 and a crystallite size of ∼480nm.

Fig. 3.
(0.2MB).

Dislocation density and crystallite size as a function of the number of (a) ECAP passes and (b) HPT turns.

During ECAP processing, the dislocation density increased to ∼8.2×1013m−2 after 4 passes and then increased further to ∼4.0×1014m−2 after 16 passes, followed by a decrease to a value of ∼2.2×1014m−2 after 24 passes, as shown in Fig. 3(a). The crystallite size, on the other hand, decreased from ∼480nm to ∼172nm after 4 passes and ∼111nm after 16 passes then increased to ∼134nm after 24 passes as in Fig. 3(a).

During HPT processing, the dislocation density increased to ∼5.6×1013m−2 after 1/2 turn followed by a decrease to values of ∼3.2×1013m−2 and ∼2.8×1013m−2 after 1 and 10 turns, respectively, as shown in Fig. 3(b). The crystallite size decreased to ∼132nm after 1/2 turn then increased to ∼154nm and ∼179nm after 1 and 10 turns, respectively.

3.3Microhardness measurements

For the ECAP specimens, the microhardness measurements were recorded along the longitudinal axes of the billets. It is readily apparent from Fig. 4(a) that the average hardness increased significantly to ∼110Hv after 4 ECAP passes in comparison to ∼41Hv measured for the annealed condition. The average hardness further increased to ∼120Hv after 16 passes. This was followed by a drop after 24 passes to ∼112Hv.

Fig. 4.
(0.26MB).

Vickers microhardness measurements for oxygen-free Cu (a) along the longitudinal axis of the ECAP billet and (b) along the radius of the HPT discs.

The microhardness measurements were recorded along the 5mm radius of the discs processed by HPT as shown in Fig. 4(b). The variation of microhardness measurements between the centre and edge positions in the HPT discs was neglected in this study. It follows from Fig. 4(b) that a substantial increase in the average hardness was recorded after 1/2 turn during HPT processing. The average hardness increased to ∼133Hv after 1/2 turn then saturated at a lower value of ∼127Hv after 10 turns.

3.4Tensile properties

Fig. 5 shows engineering stress-engineering strain curves truncated to the peak stress, which demonstrate the strain hardening behaviour for (a) ECAP and (b) HPT samples when tested at 1.0×10−4s−1. The values of the tensile properties are summarized in Table 1. As can be seen from Fig. 5(a), the mechanical strength increased for up to 16 passes followed by a drop after 24 passes. This drop in strength was associated with an increase in the uniform elongation from 2.1% to 3.6%. The same trend was observed during HPT, however, it occurred at a faster rate. The mechanical strength decreased after 1 turn while the uniform elongation increased with further strain, as in Fig. 5(b).

Fig. 5.
(0.18MB).

Truncated engineering stress – engineering strain curves demonstrating the strain hardening behaviour for (a) ECAP and (b) HPT specimens.

Table 1.

Values of yield stress (YS), ultimate tensile stress (UTS) and uniform elongation (UEL%) for Cu specimens subjected to HPT and ECAP.

Process condition  YS (MPa)  UTS (MPa)  UEL (%) 
4 passes  382  404  2.1 
16 passes  376  426  2.4 
24 passes  358  415  3.6 
½ turn  474  512  2.0 
1 turn  464  497  2.6 
10 turns  444  487  4.0 
4Discussion

Earlier research demonstrated that grain boundary sliding is the reason for the enhancement of both strength and ductility due to the increase of strain rate sensitivity [43]. The paradox of strength and ductility was first observed by processing pure Cu by ECAP for up to 16 passes and pure Ti by HPT for 5 turns at room temperature. However, this study has not focused on the intrinsic properties of the material. In the present study, oxygen-free Cu was processed by ECAP and HPT and then the crystallite size and dislocation density were examined and correlated with the mechanical properties such as microhardness and tensile properties. It was found that the hardness and strength dropped slightly at a certain strain, while uniform plastic elongation of the specimens was enhanced by processing using either ECAP or HPT.

UFG materials normally show an onset of early necking when pulled in tension, due to their low work hardening rate and low strain rate sensitivity after SPD processing, which leads to a fall in their uniform plastic deformation. The high rate of work hardening is caused by the accumulation of dislocations within the grains; however, in UFG metals the grains are small which makes dislocation storage more difficult [58]. Dislocations are emitted and absorbed at the grain boundaries instead of accumulating within the grain interior. Thus, most UFG metals exhibit limited ductility at room temperature [66]. The low hardening capacity in nanostructured Cu produced by ECAP was discussed in an earlier study and three strategies were suggested to prolong the uniform tensile deformation [67]. These strategies included creating a microstructure with a bimodal grain size distribution, deforming the material at a low temperature and/or high strain rates and increasing the strain rate sensitivity [67].

The strain rate sensitivity is a crucial factor that is considered when evaluating the strain hardening of the material. According to Hart's criterion, increasing the strain rate sensitivity can delay the onset of localized deformation and prolong the ductility of the material [68]. UFG materials tend to exhibit higher strain rate sensitivity at low temperatures in comparison to their coarse-grained counterparts [69,70]. The strain rate sensitivity for polycrystalline materials is defined as [68]:

where σ is the true stress and ε˙ the strain rate. In practice, it is well documented that a large value of m in the presence of small grains and a large fraction of HAGBs can trigger grain boundary sliding [71]. It has been reported that grain boundary sliding occurred in UFG Cu during the ECAP process at room temperature [12,43]. The occurrence of grain boundary sliding generally requires a high homologous temperature (∼0.5Tm) [72,73], however, it has been reported that grain boundary sliding also occurs at lower temperatures during ECAP and HPT processing [12,74–76]. The grain boundaries are in a nonequilibrium state during processing by ECAP or HPT due to the surplus of extrinsic dislocations which are not geometrically necessary [7] and grain boundary sliding is therefore facilitated by their movement [77].

Thus, in order to investigate the influence of the strain rate sensitivity on the ductility of the Cu specimens, another set of samples were pulled in tension using a uniaxial tensile test at a strain rate of 1.0×10−3s−1 in order to calculate the strain rate sensitivity using Eq. (3). Initially, the m value of the annealed specimen was ∼0.0083. The strain rate sensitivity increases with increasing number of ECAP passes as well as increasing HPT turns. The calculated m values of the ECAP specimens were ∼0.00941, ∼0.0193 and ∼0.0216 for 4, 16 and 24 passes, respectively, whereas, for the HPT specimens they were ∼0.036, ∼0.039 and ∼0.045 for 1/2, 1 and 10 turns, respectively. This is in agreement with an earlier report showing an increase in the strain rate sensitivity from 0.007 to 0.023 after processing pure Cu by ECAP from 1 to 12 passes [78]. There may be a small contribution towards the delay of the onset of early necking by increasing the strain rate sensitivity with ECAP passes and HPT turns, however the value is very small to facilitate grain boundary sliding.

The work hardening rate of the material is another factor contributing to the onset of the localized deformation during tensile testing as given by the Considère criterion [79,80]:

where σ is the true stress, ¿ is the true strain and ε˙ is the strain rate. Plastic instability occurs when the work hardening rate ∂σ∂ε becomes equal to or less than the flow stress during tension. Generally, the grains become smaller with increasing strain during SPD processing, which decreases the dislocation storage capacity and limits the uniform deformation. Fig. 6 presents the work hardening rate as a function of the true stress for oxygen-free Cu processed by (a) ECAP and (b) HPT. It is readily apparent from Fig. 6(a) that the specimen processed by ECAP for 24 passes exhibited a higher work hardening rate than the specimen processed by 4 passes. Also, it is apparent from Fig. 6(b) that the specimen processed by HPT for 10 turns exhibited a higher work hardening rate than the specimen processed by 1/2 turn. This indicates that the work hardening rate increases with numbers of passes or turns.

Fig. 6.
(0.21MB).

Work hardening rate as a function of true stress for oxygen-free Cu deformed by (a) ECAP and (b) HPT.

As shown in Fig. 3(a) the crystallite size decreases with increasing numbers of passes and reaches a minimum value in the region of 16 passes but then increases at 24 passes. On the other hand, there is a maximum in the dislocation density for a specimen processed by 16 passes, after which the density decreases. Similar trends were observed in Fig. 3(b) during the HPT process, where the crystallite size reached a minimum after 1/2 turn then increased with further straining, whereas the maximum value of the dislocation density occurred with the specimen subjected to 1/2 turn after which a rapid drop was observed at 1 turn followed by a continuous and gradual decrease in the dislocation density up to 10 turns. These trends of crystallite size and dislocation density displayed in Fig. 3(a) for ECAP specimens and in Fig. 3(b) for HPT specimens matches the drop in microhardness values shown in Fig. 4(a and b) and are also in a good agreement with the decrease in the strength shown in Fig. 5(a and b). These observations suggest the occurrence of a recovery mechanism during ECAP and HPT processing.

In the present study, the occurrence of the recovery mechanism highly influences the work hardening rate of oxygen-free Cu. The increase of the uniform plastic deformation presented in Table 1 is due to the decrease in both strength and dislocation densities observed during both processes, as shown in Fig. 3(a and b). This decrease in the dislocation density enhances the dislocation storage and increases the ability to accommodate more dislocations and regain the capacity of work hardening. The decrease in the dislocation density with the increase in the crystallite size is attributed to the recovery mechanism occurring during the ECAP and HPT processing that increases the mean free path of dislocations, thus, increasing the work hardening rate. This is consistent with a previous study on pure Cu processed by ECAP for 16 passes at room temperature [25].

A close inspection of Fig. 3 shows that the occurrence of the recovery mechanism during HPT is faster than for ECAP. This may be due to the intensive strain imposed by HPT. Accordingly, the equivalent strains for ECAP specimens are calculated using [59]:

where N is the number of passes, Φ is the internal channel angle and Ψ is the outer arc of curvature. The equivalent strains for 4, 16 and 24 passes are ∼3, ∼12 and ∼18, respectively.

The equivalent strains for HPT specimens were calculated using [81]:

where N is the number of rotations, r is the distance from the disc centre and h is the thickness of the disc. Since the strain varies across the disc, an average strain is calculated for each condition. The equivalent strains for 1/2, 1 and 10 turns are ∼7, ∼16 and ∼173, respectively. It is readily apparent that the recovery mechanism takes place at an equivalent strain exceeding ∼12. Thus, it is concluded that uniform plastic deformation can be prolonged by deforming the oxygen-free Cu beyond a certain equivalent strain.

5Summary and conclusions

  • 1.

    Dislocation density increased and crystalline size decreased for oxygen-free Cu after deformation by ECAP and HPT, however, at a certain strain the dislocation density decreased while the crystallite size increased indicating the occurrence of a recovery mechanism. The occurrence of this recovery is much faster during HPT than ECAP. This is attributed to the intensive strain imposed by HPT.

  • 2.

    Microhardness measurements are in agreement with the XRD results. The Hv values decreased after 24 ECAP passes and after 1 turn of HPT. The microhardness values increased significantly after ECAP and HPT deformation by comparison to the annealed condition.

  • 3.

    Although the strain rate sensitivity values are very small, it appears that they assist in the delay of early necking but not to the extent of facilitating a grain boundary sliding mechanism.

  • 4.

    The uniform plastic elongation of oxygen-free Cu was enhanced after ECAP and HPT processing as a result of the recovery mechanism whereby the dislocation annihilation process reduces the dislocation density in the presence of high-angle grain boundaries and also due to the increase in the crystallite size, which increases the mean free path of dislocations and restores the work-hardening ability of the material.

  • 5.

    The uniform plastic elongation is improved in oxygen-free Cu by imposing equivalent strains higher than ∼12.

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgements

This research was presented at the 3rd Pan American Materials Congress held as part of the TMS Annual Meeting in San Diego, California, on February 27–March 2, 2017. This work was supported by the European Research Council under ERC Grant Agreement No. 267464-SPDMETALS and by the Public Authority for Applied Education and Training (PAAET) in Kuwait.

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Paper was a contribution part of the 3rd Pan American Materials Congress, February 26th to March 2nd, 2017.

Corresponding author. (Yi Huang y.huang@soton.ac.uk)
Copyright © 2017. Brazilian Metallurgical, Materials and Mining Association