Journal of Materials Research and Technology Journal of Materials Research and Technology
J Mater Res Technol 2017;6:348-54 DOI: 10.1016/j.jmrt.2017.05.008
Original Article
Examining the microhardness evolution and thermal stability of an Al–Mg–Sc alloy processed by high-pressure torsion at a high temperature
Pedro Henrique R. Pereiraa,b,, , Yi Huanga, Terence G. Langdona
a Materials Research Group, Faculty of Engineering and the Environment, University of Southampton, Southampton, United Kingdom
b CAPES Foundation, Ministry of Education of Brazil, Brasília, DF, Brazil
Received 10 April 2017, Accepted 23 May 2017
Abstract

An Al–3% Mg–0.2% Sc alloy was solution treated and processed through 10 turns of high-pressure torsion (HPT) at 450K. Afterwards, the HPT-processed alloy was annealed for 1h at temperatures ranging from 423 to 773K and its mechanical properties and microstructural evolution were examined using microhardness measurements and electron backscattered diffraction (EBSD) analysis. The results demonstrate that HPT processing at an elevated temperature leads to a more uniform microhardness distribution and to an early saturation in the hardness values in the Al alloy compared with high-pressure torsion at room temperature. In addition, detailed EBSD analysis conducted on the HPT-processed samples immediately after annealing revealed that the Al–Mg–Sc alloy subjected to HPT processing at 450K exhibits superior thermal stability by comparison with the same material subjected to HPT at 300K.

Keywords
Aluminium alloys, Hall–Petch relationship, Hardness, High-pressure torsion, Severe plastic deformation, Thermal stability
1Introduction

There is an increasing interest in the development of Al alloys with improved mechanical properties in order to substitute heavier materials in vehicle components and reduce the energy consumption in the transportation industry [1]. Severe plastic deformation (SPD) procedures [2,3] are now recognised as effective methods in fabricating Al alloys with exceptionally small grain sizes within the submicrometre range [4]. Although various SPD techniques are now available, major focus has been dedicated to equal-channel angular pressing (ECAP) [5] and high-pressure torsion (HPT) [6] as these procedures require simple facilities and permit the processing of difficult-to-work materials by controlling the temperature, the imposed pressure and the effective strain rate in the work-pieces [7].

Ultrafine-grained (UFG) Al–Mg alloys exhibit improved mechanical strength by comparison with the material subjected to conventional metal forming [8–12]. The addition of magnesium in this material delays its recovery kinetics and promotes further grain refinement during processing by ECAP [8–10] and HPT [11,12]. Nevertheless, UFG Al–Mg alloys display significant grain coarsening after annealing at relatively low temperatures [13,14] and scandium is added to this material in order to enhance its microstructural stability and thereby preserve the benefits achieved during severe plastic deformation [15–18].

Excellent thermal stability and superplastic properties are documented for Al–Mg–Sc alloys processed by ECAP [19–21]. The Al–3Mg–0.2Sc alloy processed by ECAP at room temperature has an average grain size of ∼0.2μm immediately after processing and achieves a maximum elongation of ∼2580% during tensile testing at 723K at 3.3×10−3s−1[19]. By contrast, the Al–5Mg–0.2Sc–0.08Zr alloy has a grain size of ∼1μm after ECAP processing at 598K, however, this alloy exhibits superior superplastic properties with a record elongation of ∼4100% attained at 5.6×10−2s−1 at 723K [21]. This outstanding superplastic behaviour is attributed to a larger proportion of high-angle grain boundaries and to an improved microstructural stability in the material processed by ECAP at 598K.

Recent investigations have demonstrated that HPT processing at elevated temperatures leads to slightly larger grain sizes and lower hardness values in Mg alloys [22,23], stainless steel [24,25] and pure Ni [26] by comparison with HPT at ambient temperature. Conversely, these materials usually have more stable grain structures when annealed after HPT [24–26] and display higher superplastic elongations, as reported for the AZ61 [22] and the Mg–9Al alloy [23]. Accordingly, processing by HPT at high temperatures has emerged as a promising strategy to delay the grain coarsening kinetics in Al–Mg–Sc alloys and retain submicrometre grains at temperatures suitable for superplastic forming.

Therefore, the present study was initiated to evaluate the thermal stability of an Al–3Mg–0.2Sc alloy subjected to 10 turns of HPT processing at 450K by examining the mechanical properties and microstructural evolution in this UFG material after systematic annealing for 1h at temperatures ranging from 423 to 773K.

2Experimental material and procedures

The material used in this study was an Al–3% Mg–0.2% Sc (in wt.%) alloy supplied by China Rare Metal Material Corporation (Jiangxi Province, China) as forged bars with ∼130mm length and 10mm diameter. These bars were solution treated at 880±2K for 1h and quenched in water. Afterwards, discs with a thickness of ∼1mm were cut from the solution treated material and then ground to a final thickness of ∼0.8mm.

The discs were processed by HPT processing under quasi-constrained conditions [27,28] at 450±5K. The elevated processing temperatures were achieved by incorporating small heating elements around the upper and lower anvils and the temperature was controlled by a thermocouple placed within the upper anvil at a position of ∼10mm from the HPT sample as described in earlier investigations [22,23,29]. Initially, in the compression stage of HPT processing [30], the discs were compressed within the shallow central depression of the anvils already heated at ∼450K under a nominal pressure of 6.0GPa. Thereafter, the facility was held at the processing temperature for ∼10min and finally the lower anvil was rotated at a constant rate of 1rpm for 10 turns imposing high torsional straining within the HPT sample.

Following HPT processing, discs were annealed at temperatures from 423 to 773K for 1h and cooled in air. Afterwards, both the HPT-processed material and the annealed samples were ground and polished to obtain mirror-like surfaces and hardness measurements were then taken at the middle-section of the samples using the same procedure as in other studies [31–33]. Hardness values were evaluated along the diameters of the discs at positions separated by 0.3mm such that the microhardness in each position was calculated as the average of the measurements taken from four indentations separated by 0.15mm. Furthermore, the area-weighted average microhardness of the Al–3Mg–0.2Sc alloy was estimated using the hardness measurements recorded along the diameters of the discs and noting that the hardness values near the edges of the sample are associated with larger surface areas. The microhardness values were recorded using an FM300 microhardness tester equipped with a Vickers indenter under a load of 200gf and a dwell time of 15s.

The grain structures of the UFG metal were analysed by scanning electron microscopy (SEM) and electron backscattered diffraction (EBSD). Discs were ground, polished using 1μm diamond paste and 0.06μm colloidal silica and thereafter etched using a solution of 5% HBF4 dissolved in water. The microstructure of the Al alloy was examined using a JSM6500F thermal field emission scanning electron microscope and the average grain size, d, was estimated using the linear intercept method. EBSD patterns were collected for the materials annealed at 573, 623 and 773K using step sizes as small as 0.06μm. A cleaning procedure, including grain dilatation, was performed after data collection such that the total number of modified points was <20%. High-angle grain boundaries (HAGBs) were defined as having misorientation differences between adjacent points higher than 15° and low-angle grain boundaries (LAGBs) had misorientations within the range from 2° to 15°.

3Experimental results

Fig. 1 shows the variation of the Vickers microhardness along the diameter of the Al–3Mg–0.2Sc discs processed through 10 turns of HPT at 450K and further annealed at different temperatures. It is readily noted from these plots that there is a significant increase in the microhardness of the solution treated material after 10 turns of HPT at an elevated temperature. The Vickers microhardness of the unprocessed alloy is ∼58Hv, whereas hardness values as high as ∼185Hv were detected in the HPT-processed sample as noted in the hardness measurements denoted by open circles. Furthermore, although higher hardness values were attained in the same alloy after HPT at 300K [32], the metal processed by HPT at 450K displays a more uniform microhardness distribution with hardness values varying from ∼160 to ∼185Hv at the centre and near to the periphery of the disc, respectively.

Fig. 1.
(0.2MB).

Variation of the average microhardness recorded at the middle-section position with distance from the centre of the Al–3Mg–0.2Sc discs after processing by HPT at 450K and subsequent annealing.

It follows from Fig. 1 that post-HPT annealing at 473K leads to a slight decrease in the hardness recorded along the diameter of the Al–3Mg–0.2Sc sample, but nevertheless the overall shape of the microhardness distribution remains unchanged. In addition, the microhardness distribution in the HPT-processed metal becomes reasonably homogeneous after annealing at T>473K, although the hardness values continue decreasing with increasing temperatures.

Fig. 2 displays the variation of the area-weighted average microhardness with the annealing temperature for the Al–3Mg–0.2Sc alloy processed through 10 turns of HPT at room temperature [33] or at 450K and further annealed for 1h. The results demonstrate that the average microhardness of the material examined in this study is ∼180Hv after 10 turns of HPT at 450K. Furthermore, there is a minor reduction in the average microhardness after annealing at T473K wherein the Vickers microhardness is ∼160Hv. By contrast, it is clearly evident that annealing at temperatures within the interval from 523 to 623K leads to a more significant decrease in the hardness values which continue to decline but at lower rates for T673K.

Fig. 2.
(0.12MB).

Average microhardness as a function of annealing temperature for the Al–3Mg–0.2Sc alloy processed by 10 turns of HPT at 450K and annealed for 1h at different temperatures: for comparison purposes, additional datum points are included for the Al–3Mg–0.2Sc alloy processed by HPT at 300K and further annealed at identical conditions [33].

It is also apparent from Fig. 2 that the Al–3Mg–0.2Sc alloy processed through 10 turns of HPT at 300K exhibits higher hardness values immediately after processing compared with the same alloy processed by HPT at 450K. Nevertheless, it was consistently verified that after subsequent annealing at T473K the material processed by HPT at an elevated temperature displayed superior microhardness as a consequence of the more rapid softening kinetics for the material deformed at room temperature.

The grain boundaries of the Al–3Mg–0.2Sc alloy were revealed after etching using a solution of 5% HBF4 dissolved in water. Afterwards, various SEM images were taken in this material at positions located at distances of ∼3mm from the centres of the discs in order to estimate the average grain size after HPT at 450K and subsequent annealing at different temperatures, as depicted in Fig. 3. For comparison purposes, additional datum points are included for the same material processed by HPT at ambient temperature and further annealed at equivalent conditions [33].

Fig. 3.
(0.11MB).

Grain size as a function of annealing temperature for the Al–3Mg–0.2Sc alloy processed by 10 turns of HPT at 450K and further annealed for 1h at temperatures from 423 to 773K: for comparison purposes, additional datum points are included for the Al–3Mg–0.2Sc alloy processed by HPT at 300K and further annealed at identical conditions [33].

Fig. 3 reveals that the grain size of the Al alloy processed through 10 turns of HPT at 450K is ∼0.15μm. There is only a limited grain coarsening after heat treatment at T523K, as d<0.23μm even after annealing for 1h. Conversely, significant grain growth occurs during annealing at temperatures ranging from 573 to 673K in which the grain sizes after annealing are ∼0.35 and 1.48μm, respectively. Furthermore, it is apparent that the coarsening rate decreases to almost negligible values for T>673K such that d<2μm after annealing at 773K for 1h.

It is interesting to note that the average grain size in the Al–3Mg–0.2Sc alloy subjected to HPT at ambient temperature was only slightly inferior to the values measured in the same metal immediately after HPT at 450K. However, the experimental data displayed in Fig. 3 demonstrate that the grain structures developed in the Al alloy during HPT processing at 300K are less stable and significant grain coarsening was noted after annealing at T>423K.

Fig. 4 shows typical OIM images of the Al–3Mg–0.2Sc alloy after HPT processing at 450K and subsequent annealing for 1h, at (a) 573, (b) 623 and (c) 773K. The grains in these images were coloured using a rainbow spectrum such that small grains are blue and large structures with HAGBs are filled with shades of red. It is clearly demonstrated in Fig. 4(a) that the HPT-processed metal has a fairly homogeneous distribution of UFG structures after annealing at 573K for 1h. In addition, most of these structures are formed by HAGBs as subgrains are rarely visible in the OIM image.

Fig. 4.
(1MB).

Grain size coloured images of the grain structures of the Al–3Mg–0.2Sc alloy processed by 10 turns of HPT at 450K and further annealed for 1h at (a) 573, (b) 623 and (c) 773K.

By contrast, it is readily apparent from Fig. 4(b) that there is the onset of a bi-modal distribution of grains in the Al–Mg–Sc alloy annealed at 623K for 1h. The ultrafine grains coloured in shades of blue are uniformly distributed and represent more than 70% of the area fraction of the examined microstructure. A more detailed inspection of Fig. 4(b) revealed the grain structures were essentially equiaxed and substructures developed within some of the larger grains during annealing at 623K. In addition, the average grain size and the amount of LAGBs in the Al alloy further increased after heat treatment at 773K for 1h as observed in Fig. 4(c).

4Discussion

This investigation reveals that processing through 10 turns of HPT at 450K leads to a lower average microhardness (∼180Hv) and slightly larger grain sizes (∼0.15μm) in the Al–3Mg–0.2Sc alloy compared with HPT at room temperature [33]. These results are consistent with earlier studies that examined the effect of the processing temperature on the restoration mechanisms during high-pressure torsion [26,34]. Accordingly, lower hardness values and larger grain sizes are attained after HPT processing at increasing homologous temperatures and this is attributed to a more significant contribution of dynamic softening mechanisms such as recovery and recrystallization during severe plastic deformation, as observed in pure nickel [26], palladium and copper [34].

The plots depicted in Figs. 3 and 4 consistently confirm that the microstructural stability of Al–Mg–Sc alloys processed by HPT is notably improved by conducting this procedure at an elevated temperature. The Al–3Mg–0.2Sc alloy subjected to 10 turns of HPT at 300K and annealed at 623K exhibits an inhomogeneous microstructure formed by the few remaining fine grains and a large fraction of coarse grain structure with an average size of >10μm [33], which is similar to the duplex structures reported for pure nickel processed by HPT at 300K and further annealed at 448K for 1h [26]. By contrast, the Al alloy processed by HPT at 450K displays a more uniform distribution of grains after annealing at 623K. Although, there are a few grains with an average size of >2μm, these HAGBs structures represent less than 10% of the total amount of grains observed in Fig. 4(b).

The temperature rise of the Al–3Mg–0.2Sc alloy after 10 turns of HPT processing at ambient temperature was calculated as ∼30K using an empirical relationship [35,36]. By considering this increase in the temperature of the sample, a maximum homologous temperature, T/Tm, of ∼0.34 was estimated for the material processed at room temperature whereas T/Tm is ∼0.51 during HPT processing at 450K.

Dynamic recovery becomes more important and the rate of annihilation of dislocations increases during plastic straining at higher homologous temperatures. Accordingly, the strain energy stored in the form of microstructural defects is diminished in Al–Mg–Sc alloys processed by HPT at elevated temperatures as indicated by differential scanning calorimetry analysis conducted on an Al–5Mg–0.3Sc–0.08Zr alloy processed by HPT at different temperatures [37]. Therefore, as the strain energy provides the driving force for recrystallization, higher annealing temperatures are needed to start this phenomenon in the Al–3Mg–0.2Sc alloy processed by HPT at 450K by comparison with the alloy processed at ambient temperature.

The grain boundary strengthening during plastic straining at low homologous temperatures may be expressed in terms of the Hall–Petch relationship considering the Vickers microhardness, H, and the average grain size, d, as follows [38]:

where H0 and kH are material constants.

There are numerous studies discussing the validity of the Hall–Petch relationship for different Al–Mg–Sc alloys after SPD processing [18,33,39,40]. In order to examine the range of grain sizes in which this relationship remains accurate and assess the effect of the Mg content on the mechanical properties of Al–Mg–Sc alloys, the Vickers microhardness was plotted as a function of d−1/2 for the Al–3Mg–0.2Sc alloy processed by HPT at 450K as well as for similar Al–Mg–Sc alloys processed by HPT at ambient temperatures [33,37] or ECAP [17,18,20,41], as depicted in Fig. 5.

Fig. 5.
(0.16MB).

Plot of Vickers microhardness as a function of d−1/2 for the Al–3Mg–0.2Sc alloy processed by HPT at 450K and subsequent annealing: Additional datum points are included for similar Al–Mg–Sc alloys processed by HPT [33,37] or ECAP [17,18,20,41].

It is readily apparent from Fig. 5 that there is a linear relationship between H and d−1/2 for the Al–3Mg–0.2Sc alloy processed by HPT at an elevated temperature. Furthermore, all the experimental datum points referring to Al–Mg–Sc alloys with ∼3% Mg in weight lie close to a single line with H0=28.7Hv and kH0.056Hvm−1/2.

For comparison purposes, additional data was included for an Al–Mg–Sc–Zr alloy with higher Mg contents [18,20,41]. This Al alloy was subjected to ECAP at elevated temperatures such that grain sizes as high as ∼1μm were obtained after ECAP [18]. It is clearly noted in Fig. 5 that the Al–Mg–Sc–Zr alloys with 5 and 6% of Mg display increased hardness values by comparison with the Al–3Mg–0.2Sc alloy having the same average grain sizes. This additional hardening is a direct consequence of solid solution strengthening caused by the addition of Mg in solid solution. Nevertheless, the Hall–Petch relationship remains accurate for the ECAP-processed material with d−1/2<1000m−1/2 and kH0.056Hvm−1/2 for Al–Mg–Sc alloys with higher Mg contents.

Fig. 5 also reveals an apparent breakdown in the Hall–Petch relationship in Al–Mg–Sc alloys with d<∼50nm. This is consistent with a study wherein the grain boundary strengthening was accurately predicted using this relationship for Al–Mg–Sc alloys having grain sizes down to ∼94nm [40]. By contrast, grain refinement provides no further increase in the yield strength in pure aluminium and Al–Mg alloys without Sc additions for grain sizes smaller than ∼400 and ∼150nm, respectively.

The breakdown in the Hall–Petch relationship in nanostructured metals is usually attributed to the action of different deformation mechanisms such as dislocation emission at grain boundaries and the increasing contribution of grain boundary sliding [42]. Accordingly, the extended validity of the Hall–Petch relationship in Al–Mg–Sc alloys is associated with the segregation of Mg at grain boundaries and triple junctions during SPD processing [39,40]. In practice, these segregations reduce the boundary energy and ultimately increase the local stress needed for the emission of extrinsic dislocations [43].

Nevertheless, it is interesting to note that Mg segregation is observed in the Al–5.7Mg–0.32Sc alloy processed by HPT [44] at similar conditions as in the experimental datum points displaying no further increase in the hardness values with decreasing grain sizes. Therefore, a more comprehensive study is needed to determine whether the breakdown occurs because grain boundary sliding becomes active at low deformation temperatures or the applied stress is sufficiently high to promote the emission and annihilation of dislocations at grain boundaries in the Al–Mg–Sc alloy with average grain sizes <50nm.

Severe plastic deformation through HPT processing at room temperature promotes further strain hardening and grain refinement in Al–Mg–Sc alloys [32,37,39,45–47] by comparison with ECAP [17–20,40,48]. Conversely, the elongations to failure after tensile testing at superplastic conditions are notable inferior [45–47,49] as a consequence of the advent of abnormal grain growth at T623K [33]. Although, the onset of a bi-modal distribution of grains is also observed in the Al–3Mg–0.2Sc alloy processed through 10 turns of HPT at 450K, the grains structures in this alloy are mostly formed by HAGBs and remain essentially ultrafine (d0.85μm) even after annealing at 623K. It is concluded therefore that the microstructural stability in Al–Mg–Sc alloys is significantly improved by conducting HPT processing at 450K and this permits the retention of the benefits obtained through grain refinement and the ability to achieve superior superplastic properties at high temperatures.

5Summary and conclusions

  • 1.

    An Al–3% Mg–0.2% Sc alloy was processed through 10 turns of HPT at 450K to produce a grain size of ∼150nm. Thereafter, annealing was conducted for 1h at temperatures from 423 to 773K.

  • 2.

    A reasonably uniform hardness distribution with an average microhardness of ∼180Hv was obtained in the Al alloy after HPT. There was a continuous reduction in the hardness values during annealing at increasing temperatures, especially for T523K.

  • 3.

    The microstructural stability in Al–Mg–Sc alloys was significantly improved by conducting HPT at 450K. Dynamic recovery became more significant during HPT at high homologous temperatures and reduced the strain energy stored in the Al alloy.

  • 4.

    The microstructure is more homogeneous and the grain coarsening kinetics are delayed in the UFG metal processed by HPT at 450K by comparison with the same alloy processed by HPT at 300K. The grain structures were mostly formed by HAGBs and the average grain sizes were <2μm after annealing for 1h at 773K.

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 CAPES in Brazil and by the European Research Council under ERC Grant Agreement No. 267464-SPDMETALS.

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

Corresponding author. (Pedro Henrique R. Pereira phrp1d13@soton.ac.uk)
Copyright © 2017. Brazilian Metallurgical, Materials and Mining Association
J Mater Res Technol 2017;6:348-54 DOI: 10.1016/j.jmrt.2017.05.008