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Vol. 8. Issue 1.
Pages 630-643 (January - March 2019)
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Vol. 8. Issue 1.
Pages 630-643 (January - March 2019)
Original Article
DOI: 10.1016/j.jmrt.2018.04.016
Open Access
Dynamic strain aging behavior of an ultra-fine grained Al-Mg alloy (AA5052) processed via classical constrained groove pressing
M. Moradpoura, F. Khodabakhshib,
Corresponding author
, H. Eskandaria
a Department of Mechanical Engineering, Persian Gulf University, Bushehr 75168, Iran
b School of Metallurgical and Materials Engineering, College of Engineering, University of Tehran, P.O. Box: 11155-4563, Tehran, Iran
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Figures (9)
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Tables (3)
Table 1. Chemical composition of the utilized AA5052 alloy and comparison with the desired standard condition (wt%).
Table 2. Main mechanical properties of the initial and processed materials with using classical CGP method: yield stress (σy, MPa); ultimate tensile strength (σUTS, MPa); fracture stress (σf, MPa); elongation to failure (e, %); mean indentation micro-hardness (HV, Vickers); inhomogeneity factor (I.F., %); average cell/grain size (AGZ, nm); dislocation density (ρ, m−2).
Table 3. Contribution of the different strengthening mechanisms on the yield strength of initial annealed, wrought, and severely deformed Al-Mg alloys.
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Additional material (1)

In this study, severe plastic deformation (SPD) was employed on an Al-Mg alloy (AA5052 series) in the annealed condition with sheet form geometry by using the classical constrained-groove pressing (CGP) process up to two passes at room temperature, and accordingly an equivalent plastic strain of ∼2.32 imposed into the sheets. In relation to the microstructural features (cellular structure formation and precipitates morphology), dynamic strain aging behavior of these processed Al-Mg alloys up to different passes were evaluated by elaborating the tensile property along rolling (RD) and transverse directions (TD) of sheets in terms of anisotropy. Also, the effects of thermo-mechanical treatment in H34-temper condition on the characteristics of Al-Mg alloy was elaborated and compared with the SPD processing, as well. The results showed refinement of coarse grain structure (∼50μm) for initial annealed alloy into the ultra-fine range (400–500nm) after implementation of two CGP passes with significant homogenous enhancement of mechanical property. Hardness, yield, and tensile strengths were continuously improved up to ∼55%, 110%, and 20%, respectively, with a considerable deterioration of elongation (>90%). Notably, the strengthening mechanisms were elaborated by dislocation-based models to establish a microstructure–mechanical strength relationship.

Severe plastic deformation (SPD)
Constrained groove pressing (CGP)
Al-Mg alloy
Dynamic strain aging
Mechanical property
Full Text

As it is well established, grain refinement according to the Hall–Petch relation is a very effective method to enhance the strength of metals and alloys [1]. Materials can be categorized into three different classes based on their grain size scale as; micro-crystalline (>1μm), ultra-fine grained (100–1000nm), and nano-crystalline (<100nm) [2]. In the last two decades, improving the mechanical performance of different metals and alloys by fabrication of ultra-fine grained (UFG) and nano-structured (NS) materials have been attracted significant attentions due to their superior mechanical property and, therefore, rapid advances achieved [3]. These UFG/NS materials can be processed by means of severe plastic deformation (SPD) concept [2]. In SPD, a large amount of plastic deformation is imposed on the metal/alloy sample without any changing in its geometrical cross-sectional dimensions. Under this condition, significant increscent of dislocation density leads to the formation of dense dislocation walls and their transformation into the high-angle grain boundaries (HAGBs) [4]. SPD is mainly refined the grain structure of processed materials that can leads to significant improvements of tensile, wear, and fatigue properties [4–6]. These processed non-porous materials by SPD concept possess unique importance's including superior mechanical property as well as the appropriate dimension to perform the related physical and mechanical testing's [4,6]. Aluminum and its alloys have been broadly utilized in the various industrial applications such as transportation, aerospace and construction attributing to their high strength to weight ratio, good formability, and environmental corrosion resistance [7]. However, extending of their wide-spread application in numerous fields is restricted owing to the low hardness and wear resistance [8]. Alloying with the addition of different elements, heat treatment, and thermo-mechanical processing can be considered as the main conventional methods for improving the mechanical property of aluminum matrix by refining its microstructure [9]. However, enhancement of mechanical strength with these well-known conventional procedures is limited. Very broad researches have been performed on enhancement of the mechanical property of aluminum and its alloys by using the SPD concept and fabrication of UFG structures [2,10–13].

Several methods have been introduced for accomplishment of SPD to process the UFG or NS materials, such as; high pressure torsion (HPT) [14], mechanical alloying (MA) [15], accumulative roll bonding (ARB) [16], equal channel angular pressing (ECAP) [13], friction-stir processing (FSP) [17], constrained groove pressing (CGP) [18,19], repetitive corrugation and straightening (RCS) [20], constrained groove rolling (CGR) [21], cyclic extrusion compression (CEC) [22], twist extrusion (TE) [23], and severe torsional straining (STS) [24]. Among them, ARB, CGR, RCS, and CGP processes can be employed for SPD of sheet form metals [2,10]. Recently, classical route of CGP process, which has been proposed by Shin et al. at 2002 [19], gained many interests. In the classical route of CGP process, with alternate pressings of a sheet form material between asymmetrically corrugating and straightening dies under the plastic strain condition, repetitive shear deformation can be imparted [10]. One pass of CGP process in its classical route includes two groove pressing and two flattening stages and as a result of it, an equivalent plastic strain of around 1.16 is imposed into the processed sheet [10]. In several researches [10], this process was utilized to fabricate the UFG structure within the different metals or alloys and subsequently enhance their mechanical properties, such as; commercial pure aluminum (AA1050) [19,25], aluminum-manganese alloy (AA3003) [26], commercial pure copper [27], commercial pure nickel [28], commercial pure titanium [29], magnesium alloy (AZ31) [30], brass (Cu-Zn alloy) [31], and low carbon steel [18,32–35]. In these researches, it was found that the capability of different materials for imposing the large plastic strains with implementation of CGP process up to multi-passes are completely different [10,19,28,31,32,36,37]. For instance, copper has capability only up to three passes (equivalent plastic strain of ∼3.48) [38], while, for commercial pure aluminum the process can be continued up to even seven passes (equivalent plastic strain of ∼8.12) [39]. In the all of examined metals and alloys, an impact enhancement of mechanical strength was noticed after implementation of first CGP pass, due to drastic increscent of dislocation density and formation of a cellular structure from these dislocation tangles [18,29,30,40,41]. Thereafter, by increasing the CGP pass number more than one and imposing of higher plastic strains, the low-angle cellular structure can be altered to a high-angle one with slight size reduction and accordingly decreases the rate of mentioned improvement in the mechanical property [18,33,34,42]. Even, at higher passes some degrees of reduction in the tensile strength were reported due to occurrence of dynamic recovery (DRV) upon straining and formation of some micro-cracks [25,39,43].

Among different series of aluminum alloys, aluminum-magnesium (Al-Mg, 5XXX) series is an important class of non-heat treatable alloys which have attracted significant attentions for various functional and structural applications in aerospace and automotive industries such as lighting products, aluminum sliding, memory-disk substrates, marine engine components, gutters, cryogenic tanks, and inner body panels owing to their low cost, good recycling potential, excellent high strength to weight ratio, high fatigue strength, better corrosion resistance, and good weldability as compared to the aluminum-copper (Al-Cu, 2XXX) and aluminum-zinc (Al-Zn, 7XXX) alloys [7,8]. Moreover, the well-known serrated yielding flow behavior or Portevin-Le-Chatelier (PLC) effect that leads to occurrence of plastic instabilities in the tensile flow curve is another interesting phenomenon about these Al-Mg alloys due to the presence of Mg atoms and their locking influence action against the mobile dislocations [44,45]. In the recent studies [45–48], it was shown that by employing of SPD process on the Al-Mg alloys the most of their physical and mechanical properties such as thermal stability, work hardening, strain rate sensitivity, tensile strength and ductility as well as the fatigue strength can be improved, simultaneously. However, to date based on the authors best knowledge, there is no report on the improvement of mechanical properties for Al-Mg alloys by employing the SPD concept with using the classical route of CGP process at room temperature. Therefore, the main object of the present research was to improve the mechanical property of an AA5052 aluminum-magnesium alloy by subjecting it to CGP process at room temperature and controlling the solid-solution, grain size, and dislocation strengthening mechanisms. Also, the effects of SPD process on the serrated dynamic strain-aging flow behavior of the examined Al-Mg alloy were assessed. Finally, the microstructure–mechanical property relationship was interpreted aiming the dislocation-based micro-mechanical models to determine the contribution of different strengthening mechanisms.

2Materials and methods2.1Raw materials

In the present research, AA5052 aluminum-magnesium alloy with the chemical composition as presented in Table 1, which is in the standard range, was used as the initial raw material. The sheet of this alloy was supplied from the Arak Aluminum Company (Arak, Iran) with the thickness of 3mm and utilized into two different temper conditions of annealed and H34 (wrought). To reach the annealed condition, the purchased H34 tempered alloy was solution-treated at temperature of 500°C for 2h and subsequently water quenched.

Table 1.

Chemical composition of the utilized AA5052 alloy and comparison with the desired standard condition (wt%).

Specimens  Elements
  Al  Mg  Fe  Si  Mn  Cr  Cu  Zn 
AA5052 alloy  Base  2.26  0.295  0.0682  0.0121  0.336  0.0074  0.0131 
Standard  Base  2.2–2.8  ≤0.40  ≤0.25  ≤0.10  0.15–0.35  ≤0.10  ≤0.10 
2.2Constrained groove pressing

These annealed and wrought Al-Mg alloys were machined to the small samples with cross-sectional dimensions of 84×84mm2 before employing of CGP process. Schematic representation of classical CGP process is shown in Fig. 1. As plotted, intense plastic deformation was imposed to the sheet during CGP process in a simple shear deformation mode by repetitive pressing between the corrugation and straightening dies. The design of grooved dies was so that the distance between them was equal to the sheet thickness. Therefore, after first stage of pressing for the corrugated sheet an equivalent plastic strain of about ∼0.58 was imparted into the inclined regions placed under the deformation in plastic strain condition (see Fig. 1b). Thereafter, in the next stage of pressing another equal plastic strain was imposed to the deformed regions during straightening of the corrugated sheet between the flat dies, however, in the reverse direction. As a result, after two pressing stages the processed sheet can be divided into two regions, some parts without strain and some others with an equivalent plastic strain of ∼1.16, as shown in Fig. 1c. Next, the sheet was rotated about 180 degree around the axis perpendicular to the sheet surface to place the un-deformed regions against the inclined regions of the grooved dies and subsequently the shear deformation was applied again. With continuing the previous two corrugation and straightening stages in the new condition, an equivalent plastic strain of ∼1.16 was imparted into the un-deformed regions of the sheet, as well (see Fig. 1d and e). Therefore, after four pressing stages including two corrugations and two straightening, one pass of CGP process in its classical route as described in the literature [18,19,34] was completed and as a result an homogenous equivalent plastic strain of ∼1.16 was applied into the whole section of the processed sheet. With continuing the process up to the higher passes as two, three, four, and etc., it would be possible to impart the higher strains as 2.32, 3.48, 4.64, and so on. In this research, pressing stages were performed at room temperature under a hydraulic pressing machine with capability of one hundred ton operating at a constant cross-head rate of about 0.2mms−1. It was not possible to accomplish the classical CGP process on the AA5052 alloy in the H34 initial temper condition. Since, the sheet was broken from several regions immediately after the first stage of corrugation. For the annealed alloy, the process was performed successfully up to two passes in its classical route. At higher passes, the same problem with the wrought alloy was occurred. After examining the feasibility of CGP process implementation on the annealed and wrought Al-Mg alloys, the initial and processed samples were considered for further microstructural characterizations and mechanical testing's along the rolling (RD) and transverse directions (TD) of the sheets, as will be explained in the following sections. These samples were included the initial annealed and wrought Al-Mg alloys and CGP annealed Al-Mg alloys after one and two passes.

Fig. 1.

Schematic representation of CGP process accomplishment in its classical route.

2.3Microstructural examinations2.3.1Optical microscopy (OM) observations

Samples for optical microscopy (OM) observations were prepared from the middle parts of the initial and processed sheets. Surface sections of these extracted samples were considered for microstructural studies. After conductive mounting of these samples, standard metallographic procedure including mechanical grinding by using the SiC emery papers and polishing on diamond pasts was employed. A mixture of two chemical solutions of 0.5ml H2O-3ml HNO3-0.5ml HF-6ml HCl and 5ml H2O-12.5ml HNO3-0.5g CrO3, named as modified Poulton's reagent, was used to reveal the grain structure of the examined Al-Mg alloy. Observations were performed under an Olympus optical microscope (OM, Olympus PME3, Germany). Also, the linear intercept method was used to estimate the mean grain size of the investigated samples.

2.3.2Field emission-scanning electron microscopy (FE-SEM) analysis

The OM samples before chemical etching were considered for field emission-scanning electron microscopy (FE-SEM) analysis as well, to study the structure and morphology of precipitates inside the AA5052 alloy before and after severe plastic deformation. For this purpose, an FE-SEM instrument (FE-SEM, JEOL 7600, Japan) operating at 20keV equipped with an energy-dispersive X-ray spectroscopy (EDS) analysis detector was used.

2.3.3X-ray diffraction (XRD) analysis

To evaluate the microstructural features during CGP process with more details, X-ray diffraction (XRD) analysis technique was performed on the initial and processed samples, as well. Samples for this analysis were prepared from the middle parts of initial and processed sheets with cross-sectional dimensions of 10×20mm2. Then, their surfaces were completely polished before testing. Analysis was carried out on the angle range of 30–90° by using a Philips X-ray diffractometer equipped with a graphite monochromator using Cu Kα radiation (λ=0.1542nm). For these measurements, the step width and step time were established as 0.02° and 5s, respectively. In this method, the average grain size was determined indirectly by using the Williamson–Hall approach [18], based on the diffraction profile broadening and measuring the full-width at half maximum (FWHM). XRD profile of the annealed Al-Mg alloy was estimated as the reference state to establish the calculations.

2.3.4Transmission electron microscopy (TEM) studies

Transmission electron microscopy (TEM) analysis was carried out to directly observe and analyze the sub-grain structural refinements during CGP process. TEM samples were prepared from the middle parts of processed Al-Mg alloys with cross-sectional dimensions of 10×10mm2 and thickness of 300μm by using the electrical discharge machining (EDM) technique. Then, both sides of the extracted thin specimens were mechanically grinded and polished to reach a thickness in the range of close to ∼50μm. Afterward, these thin foils were ion-milled (JEOL, Japan) to do the hole-perforation step. Thickness around the generated holes would be in the range of 0 to 100nm, which is appropriate for TEM analysis. TEM studies were performed by using a JEOL 2000FX microscope (JEOL, Japan) operating at 200keV.

2.4Mechanical testing2.4.1Indentation hardness measurements

To assess the deformation homogeneity during CGP process and local mechanical properties enhancement, indentation Vickers micro-hardness profiles from the surface section of the initial and processed Al-Mg alloys were measured along RD and TD. These hardness measurements were accomplished by using a Vickers Bohler micro-indenter (BOHLER, Germany) and applying a load of 200g for 15s. The length and width of samples were covered at intervals of 15mm. For each point, the average value of three measurements was reported.

2.4.2Tensile testing

Tensile samples were machined along RD and TD of the initial and processed Al-Mg alloys according to the ASTM E8M Standard [49] with a gauge length of 32mm and width of 6mm by using the EDM technique. Tensile tests were conducted by using a Hounsfield Universal Tensile Testing Machine (Model H10K, Tinius Olsen, USA) at an initial strain rate of 5×10−4s−1 and room temperature. For each condition, tests were repeated at list two times. In addition to the tensile flow behavior, fracture surfaces of the tensile failed samples were studied as well under FE-SEM microscope to determine the contributing rupture mechanisms.

3Results and discussion3.1Microstructural evolutions3.1.1Deformation flow pattern and coarse grain structure

Optical microscopy images showing the microstructure of annealed Al-Mg alloy are presented in Fig. 2a–d. As seen, this started material has an equiaxed coarse grain structure with an average size of about 50μm. Also, the microstructures of wrought Al-Mg alloy in H34 temper condition are shown in Fig. 2e and f. The grain structure of this alloy is elongated along RD and refined down to ∼10μm with employing the mentioned thermo-mechanical treatment. Optical micro-graphs from the surface section of the processed Al-Mg alloy after one pass of CGP process at different magnifications are demonstrated in Fig. 3 (see Electronic Supporting Information, ESI, Fig. S1 for more information about two passes sample). Deformation flow pattern and breaking of the annealed alloy coarse grain structure after CGP process are obvious from these images. However, to measure the average grain size of the SPDed materials, optical microscopy analysis is not an appropriate technique, and as will be presented in the next sections, XRD and TEM methods were employed for this aim.

Fig. 2.

Optical micro-graphs from the grain structures of (a–d) annealed and (e and f) H34-initial temper wrought Al-Mg alloys at different magnifications.

Fig. 3.

Optical microscopy images from the different deformed regions for the surface section of CGP Al-Mg alloy after one classical pass at different magnifications; (a, d and g) 50×, (b, e and h) 100×, and (c, f and i) 200×.

3.1.2Precipitates structure and morphology

FE-SEM images from the structure of one pass and two passes processed Al-Mg alloys at different magnifications are presented in ESI Figs. S2 and S3, respectively. Some large secondary phase micro-sized particles can be observed within the AA5052 aluminum alloy matrix. With more analysis by using elemental chemical mapping analysis, as shown in ESI Figs. S4 and S5, the types of these precipitates can be studied. As seen, the most of these secondary phases are Fe-rich, and probably intermetallic phases between iron and aluminum elements. During SPD process, the initial precipitate structure of annealed Al-Mg alloy can break to the lower sizes with different morphologies by shearing effects of the induced severe plastic deformation during corrugation and straightening stages of the CGP process.

3.1.3Formation and characterization of ultra-fine grained structures3.1.3.1X-ray diffraction analysis

To monitor the microstructural evolutions with more details during SPD process, X-ray diffraction (XRD) peak profile analysis was carried out. In this method, it would be possible to indirectly estimate the average size of sub-grain/cellular structure forming during SPD process based on the broadening of XRD peaks. Then, these approximated results can be validated by direct TEM observations. XRD profiles for the annealed, wrought, and CGP Al-Mg alloys are compared in Fig. 4a. The first four (111), (200), (220), and (311) diffraction peaks are considered in calculations for evaluating the produced sub-grain/cellular structure. Variations of different peaks intensity versus the classical CGP pass number are plotted in Fig. 4b, as compared to the annealed and wrought Al-Mg alloys as reference states. As seen, some degrees of (200) and (220) preferred orientations toward the completely random and the rolled textural components are revealed in the initial annealed and wrought Al-Mg alloys, respectively. With employing of CGP process into the annealed Al-Mg alloy, intensity of all peaks is significantly decreased, although this reduction for the main peak of (200) is more considerable. This indicate tendency to the random orientation for the processed Al-Mg alloys. According to the modified Williamson-Hall approach [50,51], lattice micro-strain and crystalline structure can cause the peaks broadening which is ascribed to the sub-grain/cell size according to the following relation:

where B is the FWHM correspond to the peaks broadening, θB is the Bragg diffraction angle, ɛ is the lattice micro-strain, k is the Scherrer constant (∼0.9), λ is the wavelength of X-ray beam, and D is the mean crystalline size [51]. However, it should be noted that the peak broadening in the above equation is correspondent to the intrinsic diffraction profile, which can be measured by considering a Gaussian-Gaussian relationship for FWHM as follow [50]:
where in this equation, Bexp, B, and Bins are relating to the experimental, instrumental, and intrinsic peak broadenings, respectively. By plotting BcosθB versus sinθB, the average sub-grain/cell size can be measured by determining the intercept. The results are plotted as a function of CGP pass number and imposed strain in Fig. 4c and the main findings are summarized in Table 2. As shown, after first pass of CGP process the mean grain size of annealed Al-Mg alloy is extremely reduced from ∼50μm down to ∼660nm. Thereafter, by continuing the CGP process up to the second passes the rate of this grain structural refinement is considerably decreased and as a result an average cell size of around 500nm is attained. This finding is consistent with the general fact that the grain structural refinement is mainly considerable at the initial stages of SPD process and becomes saturated with increasing the strain [2]. Furthermore, increasing in the density of dislocations leads to diffraction profile broadening, as well. To estimate the dislocation density based on the XRD analysis, the following equations can be considered [50,52]:
where A is a calculation constant, ρ is the total dislocation density, C¯h00 is the average contrast factor for the (h00) diffraction, q is a material constant, and (hkl) is the diffraction plane Miller indices [51]. With evaluation of the XRD profiles in Fig. 4a by using these equations, dislocation density can be calculated as reported in Table 2. As seen, in relation to the grain structural refinements dislocation density is increased after the first pass of CGP process drastically from ∼1.2×1012m−2 for the annealed Al-Mg alloy up to ∼2.9×1014m−2. Thereafter, after the second passes of process it is expected more increscent up to ∼4.2×1014m−2. However, as it is well established the occurrence of dynamic recovery (DRV) which leads to dislocations annihilation can restrict the formation of cellular/UFG structures during severe plastic straining [18,34]. For wrought Al-Mg alloy, dislocation density is estimated around 3.5×1013m−2.

Fig. 4.

(a) XRD patterns for the annealed, wrought, and CGP Al-Mg alloys. (b) Evolutions of different peaks intensity as a function of CGP pass number and equivalent strain. (c) Variations of mean cell/sub-grain size during classical process versus the CGP pass number and imposed equivalent plastic strain.

Table 2.

Main mechanical properties of the initial and processed materials with using classical CGP method: yield stress (σy, MPa); ultimate tensile strength (σUTS, MPa); fracture stress (σf, MPa); elongation to failure (e, %); mean indentation micro-hardness (HV, Vickers); inhomogeneity factor (I.F., %); average cell/grain size (AGZ, nm); dislocation density (ρ, m−2).

Processing conditionsTensile propertiesHardnessMicrostructure
Materials  Initial temper  Number of CGP passes  Direction  σy  σUTS  σf  e  HV  I. FAGZ  ρ 
Al-Mg alloy  Annealed  –  RD  110.5  226.3  201.5  34.0  55.3  2.8  50,000  1.2×1012 
Al-Mg alloy  Annealed  –  TD  110.7  224.5  187.8  34.2  56.7  1.7     
Al-Mg alloy  H34  –  RD  215.4  262.7  236.9  15.4  78.4  1.2  9700  3.5×1013 
Al-Mg alloy  H34  –  TD  216.3  273.9  231.3  17.2  79.3  1.1     
Al-Mg alloy  Annealed  One  RD  226.5  251.2  202.0  8.9  81.2  2.5  664  2.9×1014 
Al-Mg alloy  Annealed  One  TD  209.7  230  207.9  5.5  78.3  2.2     
Al-Mg alloy  Annealed  Two  RD  234  260.1  197.4  7.8  85.4  2.9  495  4.2×1014 
Al-Mg alloy  Annealed  Two  TD  220.8  231.8  185.6  2.4  84.7  1.6 observations

As it is well known, observation and characterization of the UFG structures can be formed during different SPD processes, are only possible by doing the TEM analysis. TEM images from the cellular structure within the Al-Mg alloy after one and two passes of CGP process are presented in Fig. 5a–d and e–h, respectively. Intense plastic straining as caused by repetitive reverse shear straining during CGP process leads to the considerable increscent of dislocation density and formation of severely tangled dislocation structures. After one pass, formation of a combined equiaxed and elongated cellular structure during CGP process can be noticed (see Fig. 5a–d). In this processed sample, cell boundaries are mainly formed from dislocations in the tangled configuration. After employing of second passes, the microstructure became homogenous and mainly shifted toward the equiaxed sub-grains/cells, as shown in Fig. 5e–h. With more straining and imposing the higher pass numbers of CGP process, coarse elongated cells are partitioned to the equiaxed smaller ones with sharp and well-defined boundaries. Also, some portions of low-angle grain boundaries (LAGBs) are altered to the HAGBs upon restoration of equiaxed boundaries due to dislocations easy glide within the cells. The same microstructural modifications during straining by CGP process are reported by previous researchers for the other metals and alloys [10,19,21,29,41,43,53]. By using the linear intercept method and measuring the mean sub-grain/cell size from the TEM images, it can be found that the direct observations are supportive for the XRD calculations. Moreover, as shown in the high magnification TEM images of Fig. 6 some ultra-fine secondary phase Mg2Si precipitates with an average size of around 50nm are dispersed homogenously within the Al-matrix. About twenty TEM images were considered to estimate this mean size value for Mg2Si particles. This phase was not detectable in the FE-SEM images due to the detection size limit, although these ultra-fine precipitates are more effective in controlling the mechanical property of the examined Al-Mg alloy as compared to those coarse Fe-rich precipitates. Furthermore, the measured dislocation densities by XRD analysis in the previous section can be validated by considering the presented TEM images in Fig. 5a–d and e–h, as well.

Fig. 5.

TEM images showing the formation of UFG structure for the CGP Al-Mg alloy after (a–d) one and (e–h) two passes; (a, b, e and f) Bright- and (c, d, g and h) dark-field contrasts.

Fig. 6.

High magnification bright-field TEM images from the morphology and dispersion of Mg2Si precipitates within the Al-Mg alloy matrix for (a) one and (b) two passes samples.

3.2Indentation hardness

Indentation Vickers micro-hardness profiles from the surface section of initial and processed Al-Mg alloys along RD and TD are plotted in Fig. 7a and b, respectively. Also, the mean hardness values for these samples are calculated and reported in Fig. 7c and Table 2 as a function of CGP pass number at different directions. Hardness of started annealed Al-Mg alloy is varied in the range of 55–57 Vickers depending on the direction. As can be found, after one pass of classical CGP process the hardness of annealed alloy is increased significantly for both directions up to the range of 45–47%. Thereafter, the rate of this increscent is reduced in the following second passes and hardness reached to an average value of around 85 Vickers along both directions. Also, the mean hardness values for the wrought Al-Mg alloy or H34 temper condition are varied in the range of ∼78–79 Vickers along different directions, as well. The mentioned simultaneous and the same indentation hardness developments for annealed alloy along RD and TD during classical CGP process can be considered as a criterion for the negligible amount of microstructural anisotropies. Although, it seems that it is depending on the number of CGP passes, as well. To quantify the homogeneity of deformation and flow pattern during CGP of the examined Al-Mg alloy, inhomogeneity factor (I.F.) for the hardness profiles is determined by using the following relationship [18]:

where H¯, n, and Hi are the average hardness value, number of measurements, and ith one hardness value, respectively. The calculated I.F. values for the initial and processed Al-Mg alloys along RD and TD directions are reported and plotted in Table 2 and Fig. 7d, respectively. As seen after first pass of CGP process, the I.F. value increases slightly along TD and then decreases at the following second passes. While, along RD the trend is vice versa. Although, for both directions and for all processing conditions, i.e., annealed, wrought, or CGP, the magnitude of I.F. value is very small. As reported by Khodabakhshi et al. [18,33], classical CGP process was worked on the low carbon steel sheet very in-homogenous, as the degree of this non-uniformity was very high, although it was considerably decreased with increasing the CGP pass number up to four. However, one of the main results of the present research is revealed that this classical route of CGP process can undergoes very homogenous for this examined AA5052 alloy.

Fig. 7.

Vickers hardness profiles along the (a) rolling and (b) transverse directions of the annealed, wrought, and CGP alloys from the surface section. The effects of classical CGP pass number and measurement direction on the (c) mean Vickers hardness and (d) I.F. values of the annealed, wrought, and processed Al-Mg alloys.

3.3Tensile properties and dynamic strain-aging behavior

The effects of initial alloy temper condition and classical CGP process at different passes on the tensile flow behaviors of annealed Al-Mg alloy along RD and TD are presented and compared in Fig. 8a and b, respectively. The main mechanical properties obtained from these curves for different processing conditions and two measurement directions are expressed in Table 2. As shown, tempering treatment and SPD processing possess an impact influence on the dynamic strain-aging behavior of the examined Al-Mg alloy. This can be due to the increscent of dislocation density and its effects on the Portevin-Le Chatelier phenomenon [44]. The intensity of serrated yielding behavior is reduced continuously by increasing the number of CGP passes. Even, along TD after two passes, it is completely suppressed. According to the theory [44,54], unpinning and movement of dislocations during CGP process and detachment of dislocations from the Portevin-Le Chatelier atmosphere leads to disappearing/diminishing of serrated behavior in the tensile flow curves of severely deformed Al-Mg alloy. Although, material became largely strong and brittle by massive dislocations generation. Moreover, the low ductility and ductile instability in the severely deformed Al-Mg alloys could be attributed to the lack of strain hardening and occurrence of DRV mechanism which leads to decreasing of the accumulated dislocations within the sub-grains/cells by their spreading along the boundaries [2,4,54]. Variations of yield stress (YS), ultimate-tensile strength (UTS), and elongation as the main mechanical properties versus CGP pass number and imposed strain along RD and TD are demonstrated in Fig. 8c and d, respectively. As show, for both directions with increasing the CGP pass number or imposed strain the tensile strength is continuously improved up to more than two times higher, while the elongation to failure deteriorated. After first pass, the rates of these developments are considerable and thereafter in the following second passes demoted. The YS and UTS of annealed Al-Mg alloy are about 110 and 225MPa, respectively, which enhanced up to about 234 and 260MPa after two passes, at the maximum state. Although, the elongation of alloy is diminished from ∼34% down to ∼8%. Considering the tensile properties of Al-Mg alloy in its H34 temper condition, i.e., YS=215MPa, UTS=260MPa, and El=15%, more improvements in the mechanical strength were attained by using the classical CGP process.

Fig. 8.

(a and b) Effects of classical CGP process and measurement direction on the tensile flow behavior of the annealed Al-Mg alloy. The engineering stress–strain curves of Al-Mg alloy with H34-temper condition along the (a) RD and (b) TD directions plotted as well for the aim of comparison. Variations of main tensile properties versus the CGP pass number and imposed strain along the (c) rolling and (d) transverse directions.

3.4Strengthening mechanisms

To elaborate the contribution of different mechanisms including; precipitation, solid-solution, dislocation, grain boundary, and crystallographic texture on the strengthening of Al-Mg polycrystalline alloy after employing of thermo-mechanical treatment and SPD processing, a micro-mechanical model with superposition of different aspects can be employed, as below [55–57]:

where Δσgb is the strength increment due to the grain boundaries, M is a crystallographic orientation factor (often termed as the Taylor factor ∼3.06) related to texture and orientation of the material (is defined as the ratio of incremental slip on the arbitrary slip system to the magnitude of tensile or compressive incremental strain along the slip directions of 〈111〉), Δτ0 is the intrinsic strength of pure metal (∼16.7MPa), Δτss, ΔτD, and Δτppt are the solid solution, dislocation, and precipitation strengthening contributions into the total critical resolved shear stress (CRSS) of the slip planes (τtot) [56]. These strengthening mechanisms can be evaluated by considering the following equations [55,58]:
where α1 is a constant (∼2), G is the shear modulus (∼25.9 and 17GPa for Al and Mg metals, respectively), b is the Burger's vector (∼0.286nm), D is the mean grain size (Table 2), ɛ is the lattice strain due to solid-solution, c is the concentration of solute atoms (∼8.4×1028at.m−3 for the examined Al-Mg alloy), a is the lattice parameter (∼0.405 and 0.321nm for Al and Mg crystal structures, respectively), α2 is a constant (∼0.3), ρ is the dislocation density (Table 2), v is the Poisson's ratio (∼0.33), dp is the diameter of Mg2Si precipitates as measured by the TEM images in Fig. 6a and b (∼50nm), and f is the volume fraction of precipitates (∼0.4vol% for the examined Al-Mg alloy) [55,56,59]. The contribution of different strengthening mechanisms in the yield strength of the annealed, wrought, and CGP Al-Mg alloys are calculated and expressed in Table 3. As expected, grain boundary and dislocation density are being as the main and dominant factors responsible for increscent in the tensile strength of processed Al-Mg alloys with classical CGP route. A good agreement between the predicted and measured YS values can be noticed. However, for the wrought H34 aluminum alloy the difference between the simulation (∼110.7MPa) and experimental data (∼215.4MPa) is quite high. It can be mainly due to modeling assumptions and does not considering the related strengthening effects as induced by the aging phenomena.

Table 3.

Contribution of the different strengthening mechanisms on the yield strength of initial annealed, wrought, and severely deformed Al-Mg alloys.

Strengthening mechanisms  Δσgb (MPa)  Δτ0 (MPa)  Δτss (MPa)  ΔτD (MPa)  Δτppt (MPa)  Yield strength (MPa)
Al-Mg alloy condition            Predicted (σyp)  Measured (σym) 
Annealed  0.3  16.7  5.9  2.3  3.1  81.3  110.5 
Wrought-H34  1.5  16.7  5.9  12.7  3.1  110.7  215.4 
Annealed/CGP-One pass  22.3  16.7  5.9  36.5  3.1  203.5  226.5 
Annealed/CGP-Two passes  29.9  16.7  5.9  44.0  3.1  234.1  234 
3.5Fracture behavior

Fig. 9 shows the FE-SEM images from the fracture surfaces of tensile failed samples along the rolling direction. As seen, for both initial annealed and wrought Al-Mg alloys the well-developed dimples are observed over the entire fracture surfaces, which means that the fracture of these alloys are performed in a ductile behavior (see Fig. 9a–f). As grain boundaries can act as the preferred initiation sites for nucleation of voids in the micro-void coalescence mechanism, the size of dimple structure would be consistently related to the grain structure of material. For the processed UFG Al-Mg alloys after one and two passes of CGP process, very finer, rougher, and deeper dimples are appeared on the fracture surface as shown in Fig. 9g–l. Fractographic features for these two SPDed samples are exhibited a combined ductile-brittle fracture mode due to imposing of large equivalent plastic strains. According to the proposed models [60,61], by applying the tensile loading, micro-cracks can nucleate at the sub-grain/cell boundaries and interfaces between the grains and secondary phase precipitates. Thereafter, by continuing the tensile deformation they will propagate and cover the entire surface. The UFG structure can act against the crack growth and hinder the rate of its propagation by blunting or delamination mechanisms. Thereafter, based on this crack bridging and delamination model [60–62], by activation of a new slip system and subsequent cavitation, deformation can proceed in the neighbor grains which leads to the formation of UFG band. As a result, the size of dimples on the fracture surface is comparable with the UFG structure.

Fig. 9.

FE-SEM images from the fracture surfaces of (a–c) annealed, (d–f) wrought, (g–i) one pass, and (j–l) two passes CGP Al-Mg alloys.


Structural features and dynamic strain aging behavior of a commercial AA5052 Al-Mg alloy subjected to SPD at room temperature via classical route of CGP process up to different passes were comparatively investigated. The main findings can be summarized as follows:

  • Average grain sizes for the annealed and wrought Al-Mg alloys were measured by OM images around 50 and 9.7μm, respectively. Formation of a cellular structure during SPD process was characterized by TEM studies, which its mean size tremendously decreased from ∼700 to 500nm, as the number of CGP passes increased from one to two. Moreover, this grain refinement rate was retarded by the constraining effects of precipitates on the dislocations intersection, as well.

  • The average dislocation density for the CGP Al-Mg alloys was evaluated by XRD and TEM analysis, as estimated in the range of 3–4×1014m−2, which were one and two orders of magnitudes higher than the related values for the annealed and wrought Al-Mg alloys, respectively.

  • Both of thermo-mechanical and SPD treatments were found very effective in improvement the hardness and tensile property of annealed Al-Mg alloy. Although, these enhancements by using SPD processing were higher, despite of more ductility loss.

  • Both of indentation hardness and tensile strength of the CGP materials were increased dramatically as compared to the annealed base alloy. The main improvement was in the yield stress from ∼110MPa up to ∼235MPa, which was more than two times higher than those for the un-deformed alloy.

  • In relation to the different microstructural features generated along RD and TD, some degrees of in-homogeneities in the tensile and hardness responses were observed depending on the number of CGP passes. Also, the annealed Al-Mg alloy was mostly isotropic and the wrought one anisotropic.

  • Based on the proposed micro-mechanical model, the grain structural refinement and dislocation density increscent were determined as the main strengthening mechanisms in the high tensile strength of processed severely deformed Al-Mg alloy.

Conflicts of interest

The authors declare no conflicts of interest.

Appendix A
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