Deformation structures and strengthening mechanisms in an AlMgScZr alloy
Introduction
Non-age-hardenable AlMg alloys are widely used due to their good weldability, ductility, toughness, formability and high levels of corrosion resistance [1]. However, these alloys, in their annealed condition, exhibit low yield stresses (YSs), ranging from 90 to 160 MPa, depending on the Mg content [1], [2]. A magnesium concentration reaching approximately 5 wt% in aluminum alloys has been found to remain in the solid solution and has effectively increased the strength [1], [3], [4]. Minor additions of Sc and Zr to an AlMg solution lead to the formation of well-distributed, nanoscale coherent Al3(Sc,Zr) precipitates, which are thermodynamically stable. These particles are highly effective in pinning dislocations and grain boundaries, thus imparting significant strengthening and promoting microstructure stabilization [5], [6]. The strengthening mechanism of Al3(Sc,Zr) precipitates depends on their size [5]. At dispersoid sizes less than ∼25 nm, the mechanism of particle shearing dominates the particle bowing mechanism [5].
The superior properties of the AlMg materials could be further improved via the formation of an ultrafine-grained (UFG) microstructure. Severe plastic deformation (SPD) is one of the most effective approaches for this purpose [7], [8], [9], [10], [11]. Among the various SPD methods, equal-channel angular pressing (ECAP) appears to be particularly attractive due to its relative simplicity and ability to produce UFG structures in large-scale billets [7], [11], [12], [13]. The strength increase via grain refinement can be predicted by the well-known Hall-Petch (HP) relationship [7], [9], [11], [12], [14], [15], [16], [17], [18], [19], [20], [21]:σ0.2 = σo + ky d−0.5where σo is the friction stress, ky is the HP slope, and d is the crystallite size. During ECAP, continuous dynamic recrystallization (CDRX) [11], [12], [13], [15], [22], [23], [24], [25], [26], [27], [28] is the primary mechanism that results in grain refinement in aluminum alloys subjected to ECAP over a temperature interval from 250 to 400 °C. In general, CDRX includes the formation of stable, three-dimensional (3D) arrays of deformation-induced low-angle boundaries (LABs) due to dislocation rearrangement followed by their gradual transformation into high-angle boundaries (HABs) [15], [22], [23], [24], [25], [26], [27]. The new grains form due to an increase in sub-boundary misorientation resulting from the continuous accumulation of dislocations introduced by the deformation [15], [22], [23], [24], [25], [26], [27], [28]. Therefore, the average misorientation of the deformation-induced boundaries gradually increases from low to high angles as the strain increases; the deformation structure evolves from a crystallite delimited entirely by LABs to crystallites bounded partly by LABs and partly by HABs and, finally, to true UFGs bounded by HABs [23], [24], [25], [26], [27], [28].
LABs and HABs hinder dislocation glide; therefore, the number of these obstacles increases with strain that leads to increasing YSs [7], [15], [17], [18], [19], [29], [30]. The contribution of the newly formed LABs to the overall strength is strongly dependent on their misorientation. Kamikawa et al. [16] established that for pure Al subjected to SPD at room temperature, there is a critical misorientation angle that separates LABs, contributing to dislocation strengthening from boundaries that contribute to grain boundary strengthening. The LABs with misorientations of <3° provide additional strengthening through a dislocation strengthening mechanism [15], [16], [31]. However, LABs with misorientation angles greater than 3° act as conventional grain boundaries in terms of their strength contribution [16], [29], [31]. In addition, the lattice dislocation density increases by a factor ≥10 during ECAP [8], [15], [23], [24], [25], [26], [27], [28].
Presently, the contributions from the different strengthening mechanisms to the overall YS for AlMg alloys subjected to SPD are not well-understood. R.Z. Valiev et al. assumed that grain boundary strengthening in accordance with the HP law (Eq. (1)) provides the main contribution to the increment of YS due to SPD. Boundary strain-induced segregation of Mg solutes and/or precipitation of Mg-rich clusters near grain boundaries may strongly increase the efficiency of grain boundary strengthening [7], [15], [32]. From another perspective, dislocation strengthening may be the primary contributor to the overall increase in YS due to SPD in AlMg alloys [29]. Examination of the HP law (Eq. (1)) in AlMg alloys with a dislocation density ρ∼1014 m−2 without considering the contribution of the dislocation strengthening leads to an overestimation of the ky value [5], [29]. In addition, an increase in the dislocation density by thermomechanical processing [33] may significantly affect the efficiency of grain boundary strengthening in AlMg alloys. Therefore, the analysis and consideration of all strengthening mechanisms contributing to the increase in the YS due to grain refinement through ECAP is challenging. The aim of this study was to examine the effect of the deformation structure on the YS of an AlMgScZr alloy subjected to ECAP at 573 K (300 °C) in terms of the strengthening mechanisms. Specific attention was focused on establishing the role of LABs in strengthening the alloy due to SPD. The origin and nature of the effect on the deformation structure with respect to both ultimate tensile strength (UTS) and ductility will be considered in other studies.
Section snippets
Materials and methods
The alloy (i.e., 1570C Al), which had a chemical composition of Al-5.41 Mg-0.37 Mn-0.29Ti-0.2Sc-0.09Zr-0.07Fe-0.04Si (in weight %), was manufactured by direct chill casting followed by homogenization annealing at 633 K (360 °C) for 8 h and extrusion at an initial temperature of 653 K (380 °C) to produce ∼75% reduction in the cross-section, which is equal to a true strain of ∼1.3 [31]. The samples were machined from the central part of the extruded billet parallel to the extrusion direction into
Initial structure
The typical deformation structure in the initial condition is shown in Fig. 2. The relevant material microstructural characteristics are summarized in Table 1. The average dimensions of the initial grains in the extrusion direction and normal directions were ∼93 and ∼30 μm (Fig. 2a), respectively. Relatively coarse (sub)grains delimited by the LABs with misorientations typically ranging from 2 to 4° were observed. The average distance between the boundaries with misorientation ≥2° was measured
Discussion
ECAP substantially affects the grain structure and the mechanical properties. The strengthening mechanisms are discussed in this section to elucidate the key issues in the relationship between the deformation structure and the YS. Assuming that different strengthening mechanisms act independently, thus having additive contributions, the overall YS of the 1570C alloy can be expressed as follows:where is the resistance to dislocation glide within the grains for the
Conclusions
The microstructural evolution and the mechanical properties of the Al-5.4 Mg-0.2Sc-0.09Zr alloy subjected to ECAP were studied. The main conclusions from this study are as follows:
- 1.
At 300 °C, ECAP provided a +140% increase in the yield stress. The ultimate tensile strength remains nearly unchanged, and the elongation-to-failure increased slightly with an increasing number of ECAP passes.
- 2.
The increase in the yield stress was due to dislocation and boundary strengthening. Grain boundary
Acknowledgments
The financial support received from the Ministry of Education and Science, Russia, (Belgorod State University project №14.587.21.0018 (RFMEFI58715X0018)) is acknowledged. The main results were obtained by using equipment of Joint Research Center, Belgorod State University. The authors would also like to thank Anna Mogucheva for the analysis of the HP relationship and for fruitful discussions.
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