Journal of Materials Science & Technology, 2020, 52(0): 189-197 DOI: 10.1016/j.jmst.2020.04.015

Research Article

Effects of Mn content on recrystallization resistance of AA6082 aluminum alloys during post-deformation annealing

Xiaoming Qiana, Nick Parsonb, X.-Grant Chen,a,*

Department of Applied Science, University of Quebec at Chicoutimi, Saguenay G7H 2B1, Canada

Arvida Research and Development Centre, Rio Tinto Aluminum, Saguenay G7S 4K8, Canada

Corresponding authors: *.E-mail address:xgrantchen@uqac.ca(X.-G. Chen).

Received: 2019-11-23   Accepted: 2020-03-19   Online: 2020-09-15

Abstract

The microstructural evolutions under as-homogenized and as-deformed conditions and after the post-deformation annealing of AA6082 aluminum alloys with different Mn content (0.05 wt.%-1 wt.%) were studied by optical, scanning electron, and transmission electron microscopies. The results showed that the presence of a large amount of α-Al(Mn,Fe)Si dispersoids induced by Mn addition significantly improved the recrystallization resistance. In the base alloy free of Mn, static recrystallization occurred after 2 h of annealing, and grain growth commenced after 4 h of annealing, whereas in Mn-containing alloys, the recovered grain structure was well-retained after even 8 h of annealing. The alloy with 0.5% Mn exhibited the best recrystallization resistance, and a further increase of the Mn levels to 1% resulted in a gradual reduction of the recrystallization resistance, the reason for which was that recrystallization occurred only in the dispersoid-free zones (DFZs) and the increased DFZ fraction with Mn content led to an increase in the recrystallization fraction. The variation in the dispersoid number density and a coarsening of dispersoids during annealing have a limited influence on the static recrystallization in Mn-containing alloys.

Keywords: AA6082 alloys ; Mn effects ; Recrystallization resistance ; Dispersoid precipitation ; Post-deformation annealing

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Xiaoming Qian, Nick Parson, X.-Grant Chen. Effects of Mn content on recrystallization resistance of AA6082 aluminum alloys during post-deformation annealing. Journal of Materials Science & Technology[J], 2020, 52(0): 189-197 DOI:10.1016/j.jmst.2020.04.015

1. Introduction

The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg2Si precursor precipitates to attain superior mechanical properties at room temperature [1]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg2Si precursors [2]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[3], [4], [5]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [6,7]. Li et al. [8] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [9,10].

After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [11]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [12]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [13,14]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [15,16]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[17], [18], [19]].

The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [20]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[21], [22], [23], [24], [25]]. It has been well recognized that the presence of a number of fine Al3Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [21,22]. Li et al. [23] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [24] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)3Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [25] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added.

The above results on the dispersoid effect on the recrystallization resistance in aluminum alloys appear to be somewhat spread out and inconsistent owing to the complication of the recrystallization mechanism, which is attributed to the many factors involved, such as the alloying element and its content, deformation and annealing conditions. To date, no systemic studies on the effect of a Mn dispersoid-forming addition on the recrystallization resistance of 6082 alloys after a hot deformation can be found in the open literature. The evolution of the deformed microstructure during post-deformation annealing owing to the presence of numerous dispersoids, namely, the development of SRV and SRX, should be better understood.

In the present study, the effect of different Mn levels on the recrystallization resistance of AA6082 aluminum alloys was investigated. Direct chill cast billets were subjected to low-temperature homogenization at 450 °C for 6 h to promote the precipitation of Mn-containing dispersoids. The samples were then hot-deformed, followed by post-deformation annealing at 500 °C for up to 8 h. The microstructural evolutions under as-homogenized and as-deformed conditions and after post-deformation annealing were studied. Quantitative microstructural analyses were conducted on the dispersoid precipitation, SRV, and SRX to better understand the effects of dispersoids on the recrystallization resistance.

2. Experimental

Four 6082 direct chill (DC) cast alloys with Mn levels of 0.05%-1.0% were prepared for the investigation. The chemical compositions are shown in Table 1. The samples were taken from DC cast billets with a diameter of 101 mm. To produce numerous dispersoids in the aluminum matrix, DC cast billets were heat-treated at 450 °C for 6 h, followed by water quenching at room temperature, prior to hot deformation.

Table 1   Chemical composition (wt.%) of the experimental alloys.

AlloyMgSiFeMnAl
Base0.7910.180.05Bal.
0.5 Mn0.831.010.220.50Bal.
0.75 Mn0.841.020.230.72Bal.
1 Mn0.811.020.240.99Bal.

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After heat treatment, specimens with diameter 10 mm and length 15 mm were machined for compression tests, which were conducted in a Gleeble 3800 thermomechanical testing unit. The compression tests were carried out at 400 °C with a strain rate of 0.1 s-1 and a total true strain of 0.75. The specimens were heated with a heating rate of 2 °C/s to 400 °C, and then held for 180 s until the compression starts. After hot deformation, a post-deformation annealing was conducted at 500 °C for 2, 4, and 8 h to study the recrystallization behavior.

To reveal the details of the microstructure, an optical microscope (Nikon, Eclipse ME600) and a scanning electron microscope (SEM, JEOL-6480LV) were used to examine the microstructure of polished samples, which were etched with a 0.5% HF solution for 40 s. All deformed samples were sectioned parallel to the compression axis along the centerline and prepared using the standard metallographic procedure. An electron back-scattered diffraction (EBSD) analysis under SEM was applied to examine the grain structure after hot deformation and post-deformation annealing. The step size of the EBSD analysis was set to 0.5 μm. All Euler orientation maps were used in the EBSD image analysis, and boundary misorientation angles of 2°-5°, 6°-15°, and greater than 15° were used to distinguish the subgrain and grain structure. To avoid the noises caused by the sample surface preparation, misorientation angles of below 2° were not considered in the orientation maps. A transmission electron microscope (TEM, JEM - 2100) operating at 200 kV was used to observe the details of the Mn-containing dispersoids. Twin-jet polishing with 30% nitric acid and a 70% methanol solution at 15 V and -20 °C was employed to prepare the TEM thin foils.

3. Results

3.1. Microstructure after homogenization

Fig. 1 shows typical microstructures of 0.5 Mn and 1 Mn alloys as example. As indicated, the microstructures were composed of the aluminum matrix, Fe-rich intermetallics (see the inset images) along the boundaries of dendrites and grains, and primary Mg2Si (black color in the inset images) often located together with the Fe-rich intermetallics. In the base alloy with a trace of Mn, no precipitation occurred in the aluminum matrix after heat treatment. In three alloys containing Mn, numerous precipitates were observed in the aluminum matrix and were identified as α-Al(MnFe)Si dispersoids in our previous study [6]. In the 0.5 Mn alloy, the dispersoids are distributed uniformly and only very narrow dispersoid-free zones (DFZs) are observed in the interdendritic regions (Fig. 1(a)). However, with an increase in the Mn levels, the DFZs become enlarged and more recognizable. The 1 Mn alloy possesses the highest amount of DFZs (Fig. 1(b)).

Fig. 1.

Fig. 1.   Optical micrographs of 0.5 Mn (a) and 1 Mn (b) alloys after 0.5% HF etching for 40 s. The inset images are SEM micrographs of an unetched surface.


3.2. Microstructure after hot deformation

Samples after heat treatment were compressed at 400 °C. Fig. 2 shows the microstructures after hot deformation of the four alloys. Elongated grains perpendicular to the compression direction were mainly observed in deformed structures. The intermetallics were fragmented along the boundaries of dendrites and grains. In the base alloy, there are only a few separate particles of Mg2Si in the aluminum matrix but no dispersoids. In the three alloys containing Mn, fine and dense dispersoids were observed in the aluminum matrix. DFZs are also elongated along the compression direction and are more obvious in the 0.75 Mn and 1 Mn alloys. An image analysis to quantify the dispersoid number density was conducted based on a series of SEM images. The area fraction of DFZs was measured based on a series of optical images. Fig. 3 shows the number density of dispersoids and the DFZ area fraction in the three Mn-containing alloys. The number density of the dispersoids increased with an increase in the Mn amount from 8.6 μm-2 in the 0.5 Mn alloy to 10.2 and 15.3 μm-2 in the 0.75 Mn and 1 Mn alloys, respectively. For the DFZ area fraction, the 0.5 Mn alloy is as low as 2.7%, whereas it increases significantly to 6.8% and 16.7% in the 0.75 Mn and 1 Mn alloys, respectively. The sizes of the dispersoids (the mean equivalent diameter) remained nearly unchanged (∼75 nm) in the three Mn-containing alloys.

Fig. 2.

Fig. 2.   Microstructures after hot deformation: (a) the base, (b) 0.5 Mn, (c) 0.75 Mn, (d) 1 Mn alloys, (e) an enlarged SEM image of 1 Mn alloy.


Fig. 3.

Fig. 3.   The number density of dispersoids and DFZ area fraction in the different alloys after hot deformation.


The grain structures were investigated using the EBSD technique. Fig. 4 shows all Euler orientation maps of the four experimental alloys after hot deformation. In addition to the elongated grains, numerous low- and medium-angle boundaries were observed, indicating the presence of high densities of dislocations and subgrains. The deformed microstructures of all four alloys typically showed a dynamically recovered structure without dynamic recrystallization [26]. Different densities of low-, medium-, and high-angle boundaries in the four alloys were observed, representing different DRV levels [27].

Fig. 4.

Fig. 4.   All Euler orientation maps of four experimental alloys after hot deformation: (a) the base, (b) 0.5 Mn, (c) 0.75 Mn and (d) 1 Mn alloys; the white lines indicate 2°-5°, the green lines indicate 5°-15°, and the black lines indicate >15°.


The misorientation angle boundaries were analyzed based on EBSD mapping, and the results are plotted in Fig. 5. The densities of a misorientation angle of greater than 15° in all four alloys are similar (within a range of 0.14-0.19 μm-1) because the DRV during hot deformation has a limited influence on the high-angle grain boundaries. However, the density of a misorientation angle of 2°-15° (subgrain boundaries) increased from 0.35 μm-1 in the base alloy to 0.69 μm-1 in the 1 Mn alloy. The increased density of a misorientation angle of 2°-15° indicated a decline in the DRV levels with an increase in the Mn content and was believed to be related to the presence of a large amount of dispersoids. During hot deformation, the dispersoids acted as a strong barrier to the dislocation movement and subgrain migration [6,28,29]. With an increase in the dispersoid number density in the alloys, the dispersoids exerted a strong effect on the retardation of the DRV; thus, the DRV levels became lower with an increase in the Mn content (Fig. 5).

Fig. 5.

Fig. 5.   Densities of the misorientation angle boundary of experimental alloys after hot deformation.


3.3. Microstructure evolution during post-deformation annealing

Post-deformation annealing was conducted at 500 °C for 2, 4, and 8 h. All Euler orientation maps of the annealed microstructures are shown in Fig. 6 as a function of the annealing time. In addition, specific data on the boundary densities with misorientation angles of 2°-15° and >15° for the four alloys were analyzed, the results of which are plotted in Fig. 7.

Fig. 6.

Fig. 6.   All Euler orientation maps of the experimental alloys after different annealing times; the white lines indicate boundaries of 2°-5°, light green lines indicate boundaries of 5°-15°, and black lines show boundaries of >15°.


Fig. 7.

Fig. 7.   Densities of the misorientation angle boundary of 2°-5° and over 15° in (a) the base, (b) 0.5 Mn, (c) 0.75 Mn and (d) 1 Mn alloys with different annealing time.


For the base alloy after 2 h of annealing, the substructures became better organized (Fig. 6(a)) with less subgrains compared with the condition before annealing (Fig. 4(a)). Moreover, some newly formed grains were observed at near the original grain boundaries (as indicated by the arrows in Fig. 6(a)). These new grains were featured as being free of an internal substructure, indicating that partial statistic recrystallization (SRX) occurred during annealing [30,31]. The density of the boundary between 2°-15° was reduced to 0.23 μm-1 (Fig. 7(a)) compared with that before annealing (0.35 μm-1, Fig. 5), whereas the density of a boundary greater than 15° increased slightly (from 0.14 to 0.16 μm-1) owing to the occurrence of SRX, which is attributed to the formation of some new grains with high angle boundaries. With an increase in the annealing time to 4 and 8 h, abnormal grain growth occurred, whereby the grain size reached up to several hundred micrometers and few millimeters (Fig. 6(b, c)). Accordingly, the density of the boundary between 2°-15° dropped to zero, indicating no substructure within the grains, whereas a boundary density of greater than 15° decreased to close to zero (Fig. 7(a)) owing to significant grain growth.

For the three Mn-containing alloys (0.5 Mn, 0.75 Mn and 1 Mn alloys), elongated grains perpendicular to the compression direction were still observed and numerous substructures were always present after 2-8 h of annealing (Fig. 6(d-l)). In addition, no grain growth was observed after even 8 h of annealing, suggesting a much better recrystallization resistance than that of the base alloy. In the 0.5 Mn alloy, the structure of the elongated grains was well retained (Fig. 6(a-c)). However, the densities of the boundary at between 2°-15° (0.33-0.38 μm-1, Fig. 7(b)) were considerably lower than that of the condition before annealing (0.58 μm-1, Fig. 5(b)), implying the occurrence of SRV during annealing. In addition, few newly equiaxed grains without an internal substructure can be observed at the original grain boundaries (as indicated by the arrows in Fig. 6(e-f)), suggesting the start of SRX to a limited extent. These recrystallized grains were much smaller compared with those in the base alloy after 2 h of annealing. The SRX resulted in only a small increase in the densities of the boundary at greater than 15° (0.19-0.20 μm-1, Fig. 7(b)) compared with the condition before annealing (0.18 μm-1, Fig. 5). A small number of recrystallized grains became recognizable, but remained at a low level in the 0.75 Mn alloy (see the arrows in Fig. 6(g-i)). With a further increase in Mn to 1%, a larger number of recrystallized grains appeared (see the arrows in Fig. 6(j-l)) compared to the 0.75 Mn alloy. These newly formed recrystallized grains brought about a corresponding increase in the boundary density of the misorientation angles over 15° (0.21-0.23 μm-1 in Fig. 7(c) and 0.30-0.32 μm-1 in Fig. 7(d)). Meanwhile, the densities of the boundary between 2°-15° also increased with the Mn amount (Fig. 7(c, d)) owing to the increasing dispersoids number density, which led to a stronger retardation effect on the SRV.

Regarding the three Mn dispersoid-forming alloys, it was also found that the boundary structures did not exhibit an obvious change along with the annealing time (Fig. 6(d-l)), which was also reflected by the stability of the boundary densities along with the annealing time (Fig. 7(b-d)).

A quantitative image analysis was conducted on the recrystallized grain size and fraction for 0.5 Mn, 0.75 Mn and 1 Mn alloys, the results of which are shown in Fig. 8. Both the size and volume fraction of the recrystallized grains increased with an increase in the Mn addition but remained nearly unchanged with the annealing time. For instance, after 4 h of annealing, the recrystallized grain size increased from 3.8 μm in the 0.5 Mn alloy to 4.4 and 6.5 μm in the 0.75 Mn and 1 Mn alloys, respectively, whereas the volume fraction moderately increased from 2.1% in the 0.5 Mn alloy to 5.2% in the 0.75 Mn alloy and considerably to 14.7% in the 1 Mn alloy. It should be noted that, although SRX occurred in the three Mn-containing alloys, the proportion of recrystallized grains is at a low level and the deformed microstructure after post-deformation annealing mostly maintains a recovered grain structure.

Fig. 8.

Fig. 8.   Recrystallization grain size (a) and volume fraction (b) of experimental alloys at different annealing time.


The evolution of the dispersoids in the 1 Mn alloy during annealing is shown in Fig. 9. During annealing, a gradual coarsening and dissolution of the dispersoids took place, as implied by the smaller population and larger sizes of the dispersoids with an increase in the annealing time. The α-Al(MnFe)Si dispersoids are thermally stable within a temperature range of 300-350 °C, and above these temperatures they become less stable [3,5]. The results obtained here confirm that the dispersoids during annealing at 500 °C were no longer stable. The changes in the number density of the dispersoids during annealing are shown in Fig. 10. The decrease in the number density of dispersoids with the annealing time occurred in all three alloys. However, the decline in the number density was mostly significant in the 1 Mn alloy, and became less distinct in the alloy with a lower Mn content.

Fig. 9.

Fig. 9.   SEM microimages of the dispersoid evolution in the 1 Mn alloy before annealing (a) and after annealing at 500 °C for 2 h (b), 4 h (c), and 8 h (d). Bright field TEM images are inset in (a) and (c).


Fig. 10.

Fig. 10.   Number density of dispersoids in 0.5 Mn, 0.75 Mn and 1 Mn alloys at different annealing time.


The DFZ area fraction in the 0.5 Mn, 0.75 Mn and 1 Mn alloys during annealing was also measured. The results revealed that the DFZ area fraction after annealing in the three alloys remained nearly unchanged when compared with that before annealing (Fig. 3), suggesting that a coarsening and dissolution of dispersoids occurred only in the dispersoid zone and that the DFZs maintained their initial states.

4. Discussion

The effects of Mn and its dispersoid distribution on the recrystallization resistance of 6082 aluminum alloys were studied. A large amount of dispersoids were introduced through the addition of Mn and a low-temperature homogenization treatment. The presence of numerous dispersoids significantly improved the recrystallization resistance and avoided the severe grain growth during the post-deformation annealing compared with the base alloy. For the three Mn-containing alloys, the recrystallization resistance slightly decreased with an increase in the Mn addition.

In industrial practice, an excellent recrystallization resistance is desirable because a poor recrystallization resistance may cause or a recrystallized structure or a coarse grain structure during post-deformation heat treatment. In the current study, when the dispersoids were absent in the base alloy, 2 h of annealing at 500 °C resulted in the occurrence of SRX, whereas after 4 h of annealing, abnormal grain growth was observed, resulting in an extremely coarse grain structure (Fig. 6(b, c)). After the introduction of dispersoids in 0.5 Mn, 0.75 Mn and 1 Mn alloys, the deformed microstructure was stabilized and the recovered grain structure was retained (Fig. 6(d-l)), showing an excellent recrystallization resistance under high-temperature annealing treatment.

Regarding the three dispersoid-containing alloys, the number density of the dispersoids increased with an increase in the Mn content after homogenization (Fig. 3). This increase in the number density leads to an increase in the substructure density after hot deformation (Fig. 4, Fig. 5). In theory, a high number density of dispersoids can contribute to a better recrystallization resistance during annealing owing to the strong pining ability on the grain boundary migration and the grain rotation [24,32]. However, a contrary result in the present study was observed: The 0.5 Mn alloy exhibited the best recrystallization resistance with the lowest SRX fraction, whereas the 0.75 Mn alloy possessed a higher SRX fraction and the 1 Mn alloy had even the highest (Fig. 6, Fig. 8(b)).

The decreased recrystallization resistance with an increase in the Mn content was believed to be related to the DFZs. Fig. 11 shows a bright-field TEM image of the recrystallized grains in the 1 Mn alloy after 8 h of annealing. Newly formed and recrystallized grains, featured as being free of internal substructures, can be clearly observed in the interdendrite region where large intermetallic particles are present. The location of these recrystallized grains was actually in the DFZ where nearly no dispersoids existed. However, in the neighbor regions, a large amount of dispersoids remained, representing a high density of substructures in the non-recrystallized grains. This result implies that during annealing, the newly recrystallized grains preferred to nucleate and grow at the DFZ where the pinning effect of dislocations is the weakest [33]. In addition, with the increase of the Mn content from 0.5 wt.% to 1 wt.%, the amount of intermetallic particles (Fe-rich intermetallic and primary Mg2Si) moderately increased [8] and those particles were primarily located in DFZ zones. The higher amount of large intermetallic particles in the higher Mn-containing alloys can also favor the recrystallization in the DFZ zones by the particle-stimulated nucleation mechanism in some extent [34]. Once the recrystallized grains encountered the dispersoid zone, the growth was arrested, and thus the growth of recrystallized grains was restricted in the DFZ.

Fig. 11.

Fig. 11.   </b></p> </div> </div> <br> <div class="paragraph"> <div class="content-zw-1"> <p id="C26"><a class="table-icon" style="color:#2150f9" href="#F12"; id="inline_contentFig. 12">Fig. 12</a> shows additional evidence that SRX took place in the DFZs. Comparing an enlarged EBSD map with the optical image, it can be seen that the recrystallized grain band is well matched with the DFZ shape and distribution in the matrix. Furthermore, the size of the recrystallized grains (approximately 7 μm, <a class="table-icon" style="color:#2150f9" href="#F12"; id="inline_contentFig. 12">Fig. 12</a>(a)) was very close to the PFZ width (<a class="table-icon" style="color:#2150f9" href="#F12"; id="inline_contentFig. 12">Fig. 12</a>(b)). In addition, the DFZ area fractions of individual alloys correspond well with the recrystallization fractions. As mentioned above, the DFZ area fraction increased from 2.7% in the 0.5 Mn alloy to 6.8% and 16.7% in the 0.75 Mn and 1 Mn alloys, respectively (<a class="table-icon" style="color:#2150f9" href="#F3"; id="inline_contentFig. 3">Fig. 3</a>), whereas the recrystallization fraction increased from 2.1% in the 0.5 Mn alloy to 5.2% and 14.7% in the 0.75 Mn and 1 Mn alloys (<a class="table-icon" style="color:#2150f9" href="#F8"; id="inline_contentFig. 8">Fig. 8</a>(b)), respectively.</p> </div> </div> <h3 style="position: absolute; opacity: 0; filter:Alpha(opacity=0);">Fig. 12.</h3> <div class="content-zw-img" id="F12"> <div class="content-zw-img-img figure outline_anchor" onmouseleave="likai(this);"> <img src="1005-0302-52-0-189/thumbnail/img_12.jpg" onclick="clickss(this)" onmouseover="huoqukuanduimg(this);" class="tupian"> <p class="tishi"> <a href="1005-0302-52-0-189/img_12.jpg.html" target="_blank">New window</a>| <a href="1005-0302-52-0-189/img_12.jpg.zip">Download</a>| <a href="1005-0302-52-0-189/img_12.jpg.ppt"> PPT slide</a> </p> </div> <div class="content-zw-img-shuoming"> <p class="content-zw-img-shuoming-title-cn"><b>Fig. 12.   <title/> </b></p> </div> </div> <br> <div class="paragraph"> <div class="content-zw-1"> <p id="C27"><a class="table-icon" style="color:#2150f9" href="#F13"; id="inline_contentFig. 13">Fig. 13</a> shows a schematic of how the recrystallization took place. After a hot deformation, numerous substructures were induced at both the dispersoids zones and DFZs (<a class="table-icon" style="color:#2150f9" href="#F13"; id="inline_contentFig. 13">Fig. 13</a>(a)). During post-deformation annealing, the high temperature provided a driving force for the motion of dislocations and subgrains. In the dispersoid zone, owing to the strong pinning effect of the dispersoids on the dislocations and subgrain boundaries, only SRV was able to take place. However, in the DFZs, SRX can start through a diminishment or coalescence of dislocations into the subgrains owing to the absence of dispersoids and the weakest pinning effect (<a class="table-icon" style="color:#2150f9" href="#F13"; id="inline_contentFig. 13">Fig. 13</a>(b)). In addition, within the PFZs, the region surrounding the large intermetallic particles was highly strained during hot deformation, resulting in a higher density of dislocations and subgrains compared to the other regions. During annealing, the nucleation of new grains preferred to occur near the intermetallic particles where the driving force was higher, which is known as a particle-stimulated nucleation of recrystallization [<a class="demo-basic" href="javascript:;" onmouseover="jjaxxawwhah(this,'b2')">2</a>,<a class="demo-basic" href="javascript:;" onmouseover="jjaxxawwhah(this,'b34')">34</a>,<a class="demo-basic" href="javascript:;" onmouseover="jjaxxawwhah(this,'b35')">35</a>]. As a result, DFZs acted as the preferred regions where the SRX started and propagated (<a class="table-icon" style="color:#2150f9" href="#F13"; id="inline_contentFig. 13">Fig. 13</a>(c)).</p> </div> </div> <h3 style="position: absolute; opacity: 0; filter:Alpha(opacity=0);">Fig. 13.</h3> <div class="content-zw-img" id="F13"> <div class="content-zw-img-img figure outline_anchor" onmouseleave="likai(this);"> <img src="1005-0302-52-0-189/thumbnail/img_13.jpg" onclick="clickss(this)" onmouseover="huoqukuanduimg(this);" class="tupian"> <p class="tishi"> <a href="1005-0302-52-0-189/img_13.jpg.html" target="_blank">New window</a>| <a href="1005-0302-52-0-189/img_13.jpg.zip">Download</a>| <a href="1005-0302-52-0-189/img_13.jpg.ppt"> PPT slide</a> </p> </div> <div class="content-zw-img-shuoming"> <p class="content-zw-img-shuoming-title-cn"><b>Fig. 13.   <title/> </b></p> </div> </div> <br> <div class="paragraph"> <div class="content-zw-1"> <p id="C28">The coarsening and dissolution of the dispersoids occurred during the post-deformation' annealing in all three Mn-containing alloys owing to the less thermal stability of dispersoids at 500 °C (<a class="table-icon" style="color:#2150f9" href="#F9"; id="inline_contentFig. 9">Fig. 9</a>, <a class="table-icon" style="color:#2150f9" href="#F10"; id="inline_contentFig. 10">Fig. 10</a>). However, the above results indicate that the recrystallization resistance was mainly controlled by the DFZ fraction and less influenced by the number density and coarsening of the dispersoids. It was reported that even a low density of dispersoids of 0.003 μm<sup>-2</sup> can have a significant influence on the recrystallization resistance in Al-Mg-Si alloys [<a class="demo-basic" href="javascript:;" onmouseover="jjaxxawwhah(this,'b33')">33</a>]. The dispersoid densities after 8 h of annealing at 500 °C still ranged from 3.2-5.5 μm<sup>-2</sup> in the three Mn-containing alloys (<a class="table-icon" style="color:#2150f9" href="#F10"; id="inline_contentFig. 10">Fig. 10</a>), which is probably sufficient for inhibiting recrystallization. Therefore, the distribution of dispersoids associated with the DFZ is in fact the predominant factor controlling the recrystallization resistance during the post-deformation annealing.</p> </div> </div> <div class="paragraph"> <div class="content-zw-1"> <p id="C29">It is worthwhile to note that for structural applications, the hot-deformed 6082 alloys after post-deformation annealing/solution usually undergo an artificial aging to precipitate the nanoscale β”/β'-Mg<sub>2</sub>Si phases and to achieve the adequately high strengths. The presence of a large number of dispersoids induced by Mn addition can consume a part of Si solutes, which may disfavor β”/β' precipitation. On the other side, the pre-existing dispersoids formed during homogenization could provide favorable nucleation sites for subsequent β”/β' precipitation owing to the multiple benefits of the nucleation effects between the dispersoids and Mg<sub>2</sub>Si [<a class="demo-basic" href="javascript:;" onmouseover="jjaxxawwhah(this,'b36')">36</a>]. In addition to enhanced recrystallization resistance, it would be expected that the presence of dispersoids further improves the mechanical properties of final products.</p> </div> </div> <h2 class="title-biaoti outline_anchor" level="1" id="outline_anchor_1"> 5. Conclusions </h2> <div class="paragraph"> <div class="content-zw-1"> <p id="C30">(1)Through the addition of Mn and a low-temperature homogenization treatment at 450 °C, numerous α-Al(Mn,Fe)Si dispersoids were generated in the aluminum matrix of AA6082 aluminum alloys.</p> </div> </div> <div class="paragraph"> <div class="content-zw-1"> <p id="C31">(2)During post-deformation annealing at 500 °C, the base alloy free of Mn and dispersoids exhibited the worst recrystallization resistance. After 2 h of annealing, SRX occurred, whereas after 4 h of annealing, abnormal grain growth occurred.</p> </div> </div> <div class="paragraph"> <div class="content-zw-1"> <p id="C32">(3)The presence of a large number of dispersoids greatly stabilized the deformed structure and thus significantly improved the recrystallization resistance. Even after 8 h at 500 °C annealing, the recovered grain structure was well-retained in all Mn-containing 6082 aluminum alloys.</p> </div> </div> <div class="paragraph"> <div class="content-zw-1"> <p id="C33">(4)In the Mn-containing alloys, static recrystallization took place in the dispersoid free zones and the recrystallization resistance was mainly controlled by the DFZ fraction. The increased PFZ fraction with an increase in the Mn content led to an increase in the recrystallization fraction. Among the three Mn-containing alloys, the alloy with 0.5% Mn exhibited the best recrystallization resistance owing to the minimum DFZs. With a further increase in the Mn content to 0.75% and 1%, the recrystallization resistance moderately deteriorated owing to the increase in the DFZ fraction.</p> </div> </div> <h2 class="title-biaoti outline_anchor" level="1" id="outline_anchor_1"> Acknowledgements </h2> <div class="paragraph"> <div class="content-zw-1"> <p>This work was financially supported by the Natural Sciences and Engineering Research Council of Canada (No. CRDPJ 514651-17) and Rio Tinto Aluminum through the Research Chair in the Metallurgy of Aluminum Transformation at University of Quebec at Chicoutimi.</p> </div> </div> <div class="cankaowenxian1"></div> <h2 class="title-biaoti"> <span class="outline_anchor" level="1">Reference </span> <div class="btn-group"> <button style="font-size:11px;padding:3px;" type="button" class="btn btn-info dropdown-toggle" data-toggle="dropdown"> View Option <span class="caret"></span> </button> <ul class="dropdown-menu" role="menu"> <li><a href="javascript:;" class="ref_sort" type="1">By original order</a></li> <li><a href="javascript:;" class="ref_sort" type="2">By published year</a></li> <li><a href="javascript:;" class="ref_sort" type="3">By cited within times</a></li> <li><a href="javascript:;" class="ref_sort" type="4">By Impact factor</a></li> </ul> </div> </h2> <div class="cankaowenxian"> <div id="b1" name="b1" magrefid="" class="cankaowenxian-xx"> <label class="lableq"> [1] </label> <div class="cankaowenxian-xx-title"> <mixed-citation publication-type="journal"> <person-group person-group-type="author"> <name><surname>A.</surname> <given-names>Cuniberti</given-names></name>, <name><surname>A.</surname> <given-names>Tolley</given-names></name> <name><surname>M.V.</surname> <given-names>Castro Riglos</given-names></name>, <name><surname>R.</surname> <given-names>Giovachini</given-names></name>,</person-group> <source>Mater. 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A</source>, <volume>283</volume>(<issue>2000</issue>), pp. <fpage>144</fpage>-<lpage>152</lpage> </mixed-citation> </div> <p class="cankaowenxian-xx-x" id="linked_"> <span class="jcr_"></span> <span class="cjcr_"></span> <span class="doi"> <a href="https://doi.org/10.1016/S0921-5093(00)00734-6" target="_blank">DOI</a> </span>     <a href="https://linkinghub.elsevier.com/retrieve/pii/S0921509300007346" target="_blank">URL</a>     <span class="cited" id="bd_cited_count_"></span> <span class="cited" id="cited_2032295"> <a href="javascript:;" onmouseover="jjaxxawwhah1(this,'b36')" class="bianju"> [Cited within: 1]</a> </span> </p> <div class="xiangxicankao"> <p> </p> </div> <div> <div id="article_reference_meta" style="display: none;"> <div id="article_reference_meta_b1"> <div id="article_reference_meta_b1_title" class="title_"></div> <div id="article_reference_meta_b1_citedNumber">1</div> <div id="article_reference_meta_b1_nian"></div> <div id="article_reference_meta_b1_jcr"></div> <div id="article_reference_meta_b1_cjcr"></div> <div id="article_reference_meta_b1_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b2"> <div id="article_reference_meta_b2_title" class="title_"></div> <div id="article_reference_meta_b2_citedNumber">2</div> <div id="article_reference_meta_b2_nian"></div> <div id="article_reference_meta_b2_jcr"></div> <div id="article_reference_meta_b2_cjcr"></div> <div id="article_reference_meta_b2_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... <xref ref-type="fig" rid="F13">Fig. 13</xref> shows a schematic of how the recrystallization took place. After a hot deformation, numerous substructures were induced at both the dispersoids zones and DFZs (<xref ref-type="fig" rid="F13">Fig. 13</xref>(a)). During post-deformation annealing, the high temperature provided a driving force for the motion of dislocations and subgrains. In the dispersoid zone, owing to the strong pinning effect of the dispersoids on the dislocations and subgrain boundaries, only SRV was able to take place. However, in the DFZs, SRX can start through a diminishment or coalescence of dislocations into the subgrains owing to the absence of dispersoids and the weakest pinning effect (<xref ref-type="fig" rid="F13">Fig. 13</xref>(b)). In addition, within the PFZs, the region surrounding the large intermetallic particles was highly strained during hot deformation, resulting in a higher density of dislocations and subgrains compared to the other regions. During annealing, the nucleation of new grains preferred to occur near the intermetallic particles where the driving force was higher, which is known as a particle-stimulated nucleation of recrystallization [<xref ref-type="bibr" rid="b2">2</xref>,<xref ref-type="bibr" rid="b34">34</xref>,<xref ref-type="bibr" rid="b35">35</xref>]. As a result, DFZs acted as the preferred regions where the SRX started and propagated (<xref ref-type="fig" rid="F13">Fig. 13</xref>(c)). ...</div> </div> </div> <div id="article_reference_meta_b3"> <div id="article_reference_meta_b3_title" class="title_"></div> <div id="article_reference_meta_b3_citedNumber">2</div> <div id="article_reference_meta_b3_nian"></div> <div id="article_reference_meta_b3_jcr"></div> <div id="article_reference_meta_b3_cjcr"></div> <div id="article_reference_meta_b3_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... The evolution of the dispersoids in the 1 Mn alloy during annealing is shown in <xref ref-type="fig" rid="F9">Fig. 9</xref>. During annealing, a gradual coarsening and dissolution of the dispersoids took place, as implied by the smaller population and larger sizes of the dispersoids with an increase in the annealing time. The α-Al(MnFe)Si dispersoids are thermally stable within a temperature range of 300-350 °C, and above these temperatures they become less stable [<xref ref-type="bibr" rid="b3">3</xref>,<xref ref-type="bibr" rid="b5">5</xref>]. The results obtained here confirm that the dispersoids during annealing at 500 °C were no longer stable. The changes in the number density of the dispersoids during annealing are shown in <xref ref-type="fig" rid="F10">Fig. 10</xref>. The decrease in the number density of dispersoids with the annealing time occurred in all three alloys. However, the decline in the number density was mostly significant in the 1 Mn alloy, and became less distinct in the alloy with a lower Mn content. ...</div> </div> </div> <div id="article_reference_meta_b4"> <div id="article_reference_meta_b4_title" class="title_"></div> <div id="article_reference_meta_b4_citedNumber">1</div> <div id="article_reference_meta_b4_nian"></div> <div id="article_reference_meta_b4_jcr"></div> <div id="article_reference_meta_b4_cjcr"></div> <div id="article_reference_meta_b4_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b5"> <div id="article_reference_meta_b5_title" class="title_"></div> <div id="article_reference_meta_b5_citedNumber">2</div> <div id="article_reference_meta_b5_nian"></div> <div id="article_reference_meta_b5_jcr"></div> <div id="article_reference_meta_b5_cjcr"></div> <div id="article_reference_meta_b5_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... The evolution of the dispersoids in the 1 Mn alloy during annealing is shown in <xref ref-type="fig" rid="F9">Fig. 9</xref>. During annealing, a gradual coarsening and dissolution of the dispersoids took place, as implied by the smaller population and larger sizes of the dispersoids with an increase in the annealing time. The α-Al(MnFe)Si dispersoids are thermally stable within a temperature range of 300-350 °C, and above these temperatures they become less stable [<xref ref-type="bibr" rid="b3">3</xref>,<xref ref-type="bibr" rid="b5">5</xref>]. The results obtained here confirm that the dispersoids during annealing at 500 °C were no longer stable. The changes in the number density of the dispersoids during annealing are shown in <xref ref-type="fig" rid="F10">Fig. 10</xref>. The decrease in the number density of dispersoids with the annealing time occurred in all three alloys. However, the decline in the number density was mostly significant in the 1 Mn alloy, and became less distinct in the alloy with a lower Mn content. ...</div> </div> </div> <div id="article_reference_meta_b6"> <div id="article_reference_meta_b6_title" class="title_"></div> <div id="article_reference_meta_b6_citedNumber">3</div> <div id="article_reference_meta_b6_nian">2019</div> <div id="article_reference_meta_b6_jcr"></div> <div id="article_reference_meta_b6_cjcr"></div> <div id="article_reference_meta_b6_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... <xref ref-type="fig" rid="F1">Fig. 1</xref> shows typical microstructures of 0.5 Mn and 1 Mn alloys as example. As indicated, the microstructures were composed of the aluminum matrix, Fe-rich intermetallics (see the inset images) along the boundaries of dendrites and grains, and primary Mg<sub>2</sub>Si (black color in the inset images) often located together with the Fe-rich intermetallics. In the base alloy with a trace of Mn, no precipitation occurred in the aluminum matrix after heat treatment. In three alloys containing Mn, numerous precipitates were observed in the aluminum matrix and were identified as α-Al(MnFe)Si dispersoids in our previous study [<xref ref-type="bibr" rid="b6">6</xref>]. In the 0.5 Mn alloy, the dispersoids are distributed uniformly and only very narrow dispersoid-free zones (DFZs) are observed in the interdendritic regions (<xref ref-type="fig" rid="F1">Fig. 1</xref>(a)). However, with an increase in the Mn levels, the DFZs become enlarged and more recognizable. The 1 Mn alloy possesses the highest amount of DFZs (<xref ref-type="fig" rid="F1">Fig. 1</xref>(b)). ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... The misorientation angle boundaries were analyzed based on EBSD mapping, and the results are plotted in <xref ref-type="fig" rid="F5">Fig. 5</xref>. The densities of a misorientation angle of greater than 15° in all four alloys are similar (within a range of 0.14-0.19 μm<sup>-1</sup>) because the DRV during hot deformation has a limited influence on the high-angle grain boundaries. However, the density of a misorientation angle of 2°-15° (subgrain boundaries) increased from 0.35 μm<sup>-1</sup> in the base alloy to 0.69 μm<sup>-1</sup> in the 1 Mn alloy. The increased density of a misorientation angle of 2°-15° indicated a decline in the DRV levels with an increase in the Mn content and was believed to be related to the presence of a large amount of dispersoids. During hot deformation, the dispersoids acted as a strong barrier to the dislocation movement and subgrain migration [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b28">28</xref>,<xref ref-type="bibr" rid="b29">29</xref>]. With an increase in the dispersoid number density in the alloys, the dispersoids exerted a strong effect on the retardation of the DRV; thus, the DRV levels became lower with an increase in the Mn content (<xref ref-type="fig" rid="F5">Fig. 5</xref>). ...</div> </div> </div> <div id="article_reference_meta_b7"> <div id="article_reference_meta_b7_title" class="title_"></div> <div id="article_reference_meta_b7_citedNumber">1</div> <div id="article_reference_meta_b7_nian"></div> <div id="article_reference_meta_b7_jcr"></div> <div id="article_reference_meta_b7_cjcr"></div> <div id="article_reference_meta_b7_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b8"> <div id="article_reference_meta_b8_title" class="title_"></div> <div id="article_reference_meta_b8_citedNumber">2</div> <div id="article_reference_meta_b8_nian"></div> <div id="article_reference_meta_b8_jcr"></div> <div id="article_reference_meta_b8_cjcr"></div> <div id="article_reference_meta_b8_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... The decreased recrystallization resistance with an increase in the Mn content was believed to be related to the DFZs. <xref ref-type="fig" rid="F11">Fig. 11</xref> shows a bright-field TEM image of the recrystallized grains in the 1 Mn alloy after 8 h of annealing. Newly formed and recrystallized grains, featured as being free of internal substructures, can be clearly observed in the interdendrite region where large intermetallic particles are present. The location of these recrystallized grains was actually in the DFZ where nearly no dispersoids existed. However, in the neighbor regions, a large amount of dispersoids remained, representing a high density of substructures in the non-recrystallized grains. This result implies that during annealing, the newly recrystallized grains preferred to nucleate and grow at the DFZ where the pinning effect of dislocations is the weakest [<xref ref-type="bibr" rid="b33">33</xref>]. In addition, with the increase of the Mn content from 0.5 wt.% to 1 wt.%, the amount of intermetallic particles (Fe-rich intermetallic and primary Mg<sub>2</sub>Si) moderately increased [<xref ref-type="bibr" rid="b8">8</xref>] and those particles were primarily located in DFZ zones. The higher amount of large intermetallic particles in the higher Mn-containing alloys can also favor the recrystallization in the DFZ zones by the particle-stimulated nucleation mechanism in some extent [<xref ref-type="bibr" rid="b34">34</xref>]. Once the recrystallized grains encountered the dispersoid zone, the growth was arrested, and thus the growth of recrystallized grains was restricted in the DFZ. ...</div> </div> </div> <div id="article_reference_meta_b9"> <div id="article_reference_meta_b9_title" class="title_"></div> <div id="article_reference_meta_b9_citedNumber">1</div> <div id="article_reference_meta_b9_nian">2017</div> <div id="article_reference_meta_b9_jcr"></div> <div id="article_reference_meta_b9_cjcr"></div> <div id="article_reference_meta_b9_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b10"> <div id="article_reference_meta_b10_title" class="title_"></div> <div id="article_reference_meta_b10_citedNumber">1</div> <div id="article_reference_meta_b10_nian"></div> <div id="article_reference_meta_b10_jcr"></div> <div id="article_reference_meta_b10_cjcr"></div> <div id="article_reference_meta_b10_articleCitedText"> <div class="sentence">... The traditional hardening mechanism of Al-Mg-Si 6xxx aluminum alloys is through the precipitation of fine nano-scale Mg<sub>2</sub>Si precursor precipitates to attain superior mechanical properties at room temperature [<xref ref-type="bibr" rid="b1">1</xref>]. However, for a service temperature exceeding 200 °C, the mechanical properties deteriorate rapidly owing to the coarsening and dissolution of Mg<sub>2</sub>Si precursors [<xref ref-type="bibr" rid="b2">2</xref>]. To develop alloys that can be applied at elevated temperatures, several studies have been conducted by the introduction of numerous thermally stable dispersoids in aluminum alloys [[<xref ref-type="bibr" rid="b3">3</xref>], [<xref ref-type="bibr" rid="b4">4</xref>], [<xref ref-type="bibr" rid="b5">5</xref>]]. For AA6082 aluminum alloys, one of the most common stable dispersoids encountered in the matrix is α-Al(Mn,Fe)Si, which is partially coherent and can be formed through decomposition of the supersaturated solid solution during homogenization [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b7">7</xref>]. Li et al. [<xref ref-type="bibr" rid="b8">8</xref>] reported that the precipitation of a large amount of α-Al(Mn,Fe)Si dispersoids can be promoted with the addition of Mn in 6082 alloys after a relatively low-temperature treatment at 400-450 °C, while the number density of dispersoids increases with an increase in the Mn levels. The strong effect of Mn addition on the precipitation of α-Al(Mn,Fe)Si dispersoids during high-temperature homogenization (550-580 °C) in 6082-based alloys was also reported [<xref ref-type="bibr" rid="b9">9</xref>,<xref ref-type="bibr" rid="b10">10</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b11"> <div id="article_reference_meta_b11_title" class="title_"></div> <div id="article_reference_meta_b11_citedNumber">1</div> <div id="article_reference_meta_b11_nian"></div> <div id="article_reference_meta_b11_jcr"></div> <div id="article_reference_meta_b11_cjcr"></div> <div id="article_reference_meta_b11_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b12"> <div id="article_reference_meta_b12_title" class="title_"></div> <div id="article_reference_meta_b12_citedNumber">1</div> <div id="article_reference_meta_b12_nian"></div> <div id="article_reference_meta_b12_jcr"></div> <div id="article_reference_meta_b12_cjcr"></div> <div id="article_reference_meta_b12_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b13"> <div id="article_reference_meta_b13_title" class="title_"></div> <div id="article_reference_meta_b13_citedNumber">1</div> <div id="article_reference_meta_b13_nian"></div> <div id="article_reference_meta_b13_jcr"></div> <div id="article_reference_meta_b13_cjcr"></div> <div id="article_reference_meta_b13_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b14"> <div id="article_reference_meta_b14_title" class="title_"></div> <div id="article_reference_meta_b14_citedNumber">1</div> <div id="article_reference_meta_b14_nian"></div> <div id="article_reference_meta_b14_jcr"></div> <div id="article_reference_meta_b14_cjcr"></div> <div id="article_reference_meta_b14_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b15"> <div id="article_reference_meta_b15_title" class="title_"></div> <div id="article_reference_meta_b15_citedNumber">1</div> <div id="article_reference_meta_b15_nian"></div> <div id="article_reference_meta_b15_jcr"></div> <div id="article_reference_meta_b15_cjcr"></div> <div id="article_reference_meta_b15_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b16"> <div id="article_reference_meta_b16_title" class="title_"></div> <div id="article_reference_meta_b16_citedNumber">1</div> <div id="article_reference_meta_b16_nian"></div> <div id="article_reference_meta_b16_jcr"></div> <div id="article_reference_meta_b16_cjcr"></div> <div id="article_reference_meta_b16_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b17"> <div id="article_reference_meta_b17_title" class="title_"></div> <div id="article_reference_meta_b17_citedNumber">1</div> <div id="article_reference_meta_b17_nian"></div> <div id="article_reference_meta_b17_jcr"></div> <div id="article_reference_meta_b17_cjcr"></div> <div id="article_reference_meta_b17_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b18"> <div id="article_reference_meta_b18_title" class="title_"></div> <div id="article_reference_meta_b18_citedNumber">1</div> <div id="article_reference_meta_b18_nian">2010</div> <div id="article_reference_meta_b18_jcr"></div> <div id="article_reference_meta_b18_cjcr"></div> <div id="article_reference_meta_b18_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b19"> <div id="article_reference_meta_b19_title" class="title_"></div> <div id="article_reference_meta_b19_citedNumber">1</div> <div id="article_reference_meta_b19_nian"></div> <div id="article_reference_meta_b19_jcr"></div> <div id="article_reference_meta_b19_cjcr"></div> <div id="article_reference_meta_b19_articleCitedText"> <div class="sentence">... After casting and homogenization, 6xxx aluminum alloys are usually subjected to a thermomechanical process, such as rolling or extrusion to achieve a desirable shape. The deformed structures are usually associated with a high level of internal stress and a high density of the substructures. To achieve appropriate and stable mechanical and materials properties, a post-deformation heat treatment (annealing or solution treatment) is applied [<xref ref-type="bibr" rid="b11">11</xref>]. Statistic recovery (SRV) and statistic recrystallization (SRX) can occur during the post-deformation annealing. SRV is associated with a change in the density and distribution of line defects, whereas SRX involves the nucleation and growth of new grains as well as grain boundary migration [<xref ref-type="bibr" rid="b12">12</xref>]. The control of SRX plays an important role in wrought aluminum alloys. It has been reported that the occurrence of SRX negatively influences the corrosion resistance in 2xxx alloys [<xref ref-type="bibr" rid="b13">13</xref>,<xref ref-type="bibr" rid="b14">14</xref>]. In 5xxx alloys, the work hardening effect can be kept only if the non-recrystallized structure can be maintained [<xref ref-type="bibr" rid="b15">15</xref>,<xref ref-type="bibr" rid="b16">16</xref>]. In 7xxx alloys, a recrystallized structure can cause an increasing risk of weld cracking, a decreased fracture toughness, and a detrimental effect on the corrosion resistance [[<xref ref-type="bibr" rid="b17">17</xref>], [<xref ref-type="bibr" rid="b18">18</xref>], [<xref ref-type="bibr" rid="b19">19</xref>]]. ...</div> </div> </div> <div id="article_reference_meta_b20"> <div id="article_reference_meta_b20_title" class="title_"></div> <div id="article_reference_meta_b20_citedNumber">1</div> <div id="article_reference_meta_b20_nian"></div> <div id="article_reference_meta_b20_jcr"></div> <div id="article_reference_meta_b20_cjcr"></div> <div id="article_reference_meta_b20_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> </div> </div> <div id="article_reference_meta_b21"> <div id="article_reference_meta_b21_title" class="title_"></div> <div id="article_reference_meta_b21_citedNumber">2</div> <div id="article_reference_meta_b21_nian"></div> <div id="article_reference_meta_b21_jcr"></div> <div id="article_reference_meta_b21_cjcr"></div> <div id="article_reference_meta_b21_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> </div> </div> <div id="article_reference_meta_b22"> <div id="article_reference_meta_b22_title" class="title_"></div> <div id="article_reference_meta_b22_citedNumber">2</div> <div id="article_reference_meta_b22_nian"></div> <div id="article_reference_meta_b22_jcr"></div> <div id="article_reference_meta_b22_cjcr"></div> <div id="article_reference_meta_b22_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... ,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> </div> </div> <div id="article_reference_meta_b23"> <div id="article_reference_meta_b23_title" class="title_"></div> <div id="article_reference_meta_b23_citedNumber">2</div> <div id="article_reference_meta_b23_nian"></div> <div id="article_reference_meta_b23_jcr"></div> <div id="article_reference_meta_b23_cjcr"></div> <div id="article_reference_meta_b23_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... ]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> </div> </div> <div id="article_reference_meta_b24"> <div id="article_reference_meta_b24_title" class="title_"></div> <div id="article_reference_meta_b24_citedNumber">3</div> <div id="article_reference_meta_b24_nian"></div> <div id="article_reference_meta_b24_jcr"></div> <div id="article_reference_meta_b24_cjcr"></div> <div id="article_reference_meta_b24_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... ] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... Regarding the three dispersoid-containing alloys, the number density of the dispersoids increased with an increase in the Mn content after homogenization (<xref ref-type="fig" rid="F3">Fig. 3</xref>). This increase in the number density leads to an increase in the substructure density after hot deformation (<xref ref-type="fig" rid="F4">Fig. 4</xref>, <xref ref-type="fig" rid="F5">Fig. 5</xref>). In theory, a high number density of dispersoids can contribute to a better recrystallization resistance during annealing owing to the strong pining ability on the grain boundary migration and the grain rotation [<xref ref-type="bibr" rid="b24">24</xref>,<xref ref-type="bibr" rid="b32">32</xref>]. However, a contrary result in the present study was observed: The 0.5 Mn alloy exhibited the best recrystallization resistance with the lowest SRX fraction, whereas the 0.75 Mn alloy possessed a higher SRX fraction and the 1 Mn alloy had even the highest (<xref ref-type="fig" rid="F6">Fig. 6</xref>, <xref ref-type="fig" rid="F8">Fig. 8</xref>(b)). ...</div> </div> </div> <div id="article_reference_meta_b25"> <div id="article_reference_meta_b25_title" class="title_"></div> <div id="article_reference_meta_b25_citedNumber">2</div> <div id="article_reference_meta_b25_nian"></div> <div id="article_reference_meta_b25_jcr"></div> <div id="article_reference_meta_b25_cjcr"></div> <div id="article_reference_meta_b25_articleCitedText"> <div class="sentence">... The pre-existing thermally stable dispersoids in the aluminum matrix can significantly control grain growth and retard recrystallization owing to their pinning effect on grain boundary migration [<xref ref-type="bibr" rid="b20">20</xref>]. The size, number density, and distribution of the dispersoids have a significant influence on the recrystallization resistance [[<xref ref-type="bibr" rid="b21">21</xref>], [<xref ref-type="bibr" rid="b22">22</xref>], [<xref ref-type="bibr" rid="b23">23</xref>], [<xref ref-type="bibr" rid="b24">24</xref>], [<xref ref-type="bibr" rid="b25">25</xref>]]. It has been well recognized that the presence of a number of fine Al<sub>3</sub>Zr dispersoids can significantly increase the recrystallization resistance during post-deformation annealing in 7xxx alloys [<xref ref-type="bibr" rid="b21">21</xref>,<xref ref-type="bibr" rid="b22">22</xref>]. Li et al. [<xref ref-type="bibr" rid="b23">23</xref>] studied the effects of Er and Zr on recrystallization in pure aluminum and found that Al3(Er,Zr) dispersoids can be formed during heat treatment at 400 °C for 48 h, resulting in a remarkable enhancement of the recrystallization resistance during annealing at 350-525 °C. Birol [<xref ref-type="bibr" rid="b24">24</xref>] reported that a superior recrystallization resistance of 6082 alloy can be obtained through a large population of Cr-rich Al(Cr,Mn,Fe)Si and (Al,Si)<sub>3</sub>Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... Zr dispersoids. However, the individual addition of Mn or Zr fails to offer any improvement in the recrystallization resistance in 6082 tube extrusions. Tsivoulas et al. [<xref ref-type="bibr" rid="b25">25</xref>] investigated the effects of Mn and Zr additions on the recrystallization resistance in Al-Cu-Li 2198 sheets, and found that with a constant Zr level, the recrystallization resistance was diminished with the addition of Mn, and progressively worsened with a decrease in the amount of Zr as more Mn was added. ...</div> </div> </div> <div id="article_reference_meta_b26"> <div id="article_reference_meta_b26_title" class="title_"></div> <div id="article_reference_meta_b26_citedNumber">1</div> <div id="article_reference_meta_b26_nian"></div> <div id="article_reference_meta_b26_jcr"></div> <div id="article_reference_meta_b26_cjcr"></div> <div id="article_reference_meta_b26_articleCitedText"> <div class="sentence">... The grain structures were investigated using the EBSD technique. <xref ref-type="fig" rid="F4">Fig. 4</xref> shows all Euler orientation maps of the four experimental alloys after hot deformation. In addition to the elongated grains, numerous low- and medium-angle boundaries were observed, indicating the presence of high densities of dislocations and subgrains. The deformed microstructures of all four alloys typically showed a dynamically recovered structure without dynamic recrystallization [<xref ref-type="bibr" rid="b26">26</xref>]. Different densities of low-, medium-, and high-angle boundaries in the four alloys were observed, representing different DRV levels [<xref ref-type="bibr" rid="b27">27</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b27"> <div id="article_reference_meta_b27_title" class="title_"></div> <div id="article_reference_meta_b27_citedNumber">1</div> <div id="article_reference_meta_b27_nian"></div> <div id="article_reference_meta_b27_jcr"></div> <div id="article_reference_meta_b27_cjcr"></div> <div id="article_reference_meta_b27_articleCitedText"> <div class="sentence">... The grain structures were investigated using the EBSD technique. <xref ref-type="fig" rid="F4">Fig. 4</xref> shows all Euler orientation maps of the four experimental alloys after hot deformation. In addition to the elongated grains, numerous low- and medium-angle boundaries were observed, indicating the presence of high densities of dislocations and subgrains. The deformed microstructures of all four alloys typically showed a dynamically recovered structure without dynamic recrystallization [<xref ref-type="bibr" rid="b26">26</xref>]. Different densities of low-, medium-, and high-angle boundaries in the four alloys were observed, representing different DRV levels [<xref ref-type="bibr" rid="b27">27</xref>]. ...</div> </div> </div> <div id="article_reference_meta_b28"> <div id="article_reference_meta_b28_title" class="title_"></div> <div id="article_reference_meta_b28_citedNumber">1</div> <div id="article_reference_meta_b28_nian"></div> <div id="article_reference_meta_b28_jcr"></div> <div id="article_reference_meta_b28_cjcr"></div> <div id="article_reference_meta_b28_articleCitedText"> <div class="sentence">... The misorientation angle boundaries were analyzed based on EBSD mapping, and the results are plotted in <xref ref-type="fig" rid="F5">Fig. 5</xref>. The densities of a misorientation angle of greater than 15° in all four alloys are similar (within a range of 0.14-0.19 μm<sup>-1</sup>) because the DRV during hot deformation has a limited influence on the high-angle grain boundaries. However, the density of a misorientation angle of 2°-15° (subgrain boundaries) increased from 0.35 μm<sup>-1</sup> in the base alloy to 0.69 μm<sup>-1</sup> in the 1 Mn alloy. The increased density of a misorientation angle of 2°-15° indicated a decline in the DRV levels with an increase in the Mn content and was believed to be related to the presence of a large amount of dispersoids. During hot deformation, the dispersoids acted as a strong barrier to the dislocation movement and subgrain migration [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b28">28</xref>,<xref ref-type="bibr" rid="b29">29</xref>]. With an increase in the dispersoid number density in the alloys, the dispersoids exerted a strong effect on the retardation of the DRV; thus, the DRV levels became lower with an increase in the Mn content (<xref ref-type="fig" rid="F5">Fig. 5</xref>). ...</div> </div> </div> <div id="article_reference_meta_b29"> <div id="article_reference_meta_b29_title" class="title_"></div> <div id="article_reference_meta_b29_citedNumber">1</div> <div id="article_reference_meta_b29_nian"></div> <div id="article_reference_meta_b29_jcr"></div> <div id="article_reference_meta_b29_cjcr"></div> <div id="article_reference_meta_b29_articleCitedText"> <div class="sentence">... The misorientation angle boundaries were analyzed based on EBSD mapping, and the results are plotted in <xref ref-type="fig" rid="F5">Fig. 5</xref>. The densities of a misorientation angle of greater than 15° in all four alloys are similar (within a range of 0.14-0.19 μm<sup>-1</sup>) because the DRV during hot deformation has a limited influence on the high-angle grain boundaries. However, the density of a misorientation angle of 2°-15° (subgrain boundaries) increased from 0.35 μm<sup>-1</sup> in the base alloy to 0.69 μm<sup>-1</sup> in the 1 Mn alloy. The increased density of a misorientation angle of 2°-15° indicated a decline in the DRV levels with an increase in the Mn content and was believed to be related to the presence of a large amount of dispersoids. During hot deformation, the dispersoids acted as a strong barrier to the dislocation movement and subgrain migration [<xref ref-type="bibr" rid="b6">6</xref>,<xref ref-type="bibr" rid="b28">28</xref>,<xref ref-type="bibr" rid="b29">29</xref>]. With an increase in the dispersoid number density in the alloys, the dispersoids exerted a strong effect on the retardation of the DRV; thus, the DRV levels became lower with an increase in the Mn content (<xref ref-type="fig" rid="F5">Fig. 5</xref>). ...</div> </div> </div> <div id="article_reference_meta_b30"> <div id="article_reference_meta_b30_title" class="title_"></div> <div id="article_reference_meta_b30_citedNumber">1</div> <div id="article_reference_meta_b30_nian"></div> <div id="article_reference_meta_b30_jcr"></div> <div id="article_reference_meta_b30_cjcr"></div> <div id="article_reference_meta_b30_articleCitedText"> <div class="sentence">... For the base alloy after 2 h of annealing, the substructures became better organized (<xref ref-type="fig" rid="F6">Fig. 6</xref>(a)) with less subgrains compared with the condition before annealing (<xref ref-type="fig" rid="F4">Fig. 4</xref>(a)). Moreover, some newly formed grains were observed at near the original grain boundaries (as indicated by the arrows in <xref ref-type="fig" rid="F6">Fig. 6</xref>(a)). These new grains were featured as being free of an internal substructure, indicating that partial statistic recrystallization (SRX) occurred during annealing [<xref ref-type="bibr" rid="b30">30</xref>,<xref ref-type="bibr" rid="b31">31</xref>]. The density of the boundary between 2°-15° was reduced to 0.23 μm<sup>-1</sup> (<xref ref-type="fig" rid="F7">Fig. 7</xref>(a)) compared with that before annealing (0.35 μm<sup>-1</sup>, <xref ref-type="fig" rid="F5">Fig. 5</xref>), whereas the density of a boundary greater than 15° increased slightly (from 0.14 to 0.16 μm<sup>-1</sup>) owing to the occurrence of SRX, which is attributed to the formation of some new grains with high angle boundaries. With an increase in the annealing time to 4 and 8 h, abnormal grain growth occurred, whereby the grain size reached up to several hundred micrometers and few millimeters (<xref ref-type="fig" rid="F6">Fig. 6</xref>(b, c)). Accordingly, the density of the boundary between 2°-15° dropped to zero, indicating no substructure within the grains, whereas a boundary density of greater than 15° decreased to close to zero (<xref ref-type="fig" rid="F7">Fig. 7</xref>(a)) owing to significant grain growth. ...</div> </div> </div> <div id="article_reference_meta_b31"> <div id="article_reference_meta_b31_title" class="title_"></div> <div id="article_reference_meta_b31_citedNumber">1</div> <div id="article_reference_meta_b31_nian"></div> <div id="article_reference_meta_b31_jcr"></div> <div id="article_reference_meta_b31_cjcr"></div> <div id="article_reference_meta_b31_articleCitedText"> <div class="sentence">... For the base alloy after 2 h of annealing, the substructures became better organized (<xref ref-type="fig" rid="F6">Fig. 6</xref>(a)) with less subgrains compared with the condition before annealing (<xref ref-type="fig" rid="F4">Fig. 4</xref>(a)). Moreover, some newly formed grains were observed at near the original grain boundaries (as indicated by the arrows in <xref ref-type="fig" rid="F6">Fig. 6</xref>(a)). These new grains were featured as being free of an internal substructure, indicating that partial statistic recrystallization (SRX) occurred during annealing [<xref ref-type="bibr" rid="b30">30</xref>,<xref ref-type="bibr" rid="b31">31</xref>]. The density of the boundary between 2°-15° was reduced to 0.23 μm<sup>-1</sup> (<xref ref-type="fig" rid="F7">Fig. 7</xref>(a)) compared with that before annealing (0.35 μm<sup>-1</sup>, <xref ref-type="fig" rid="F5">Fig. 5</xref>), whereas the density of a boundary greater than 15° increased slightly (from 0.14 to 0.16 μm<sup>-1</sup>) owing to the occurrence of SRX, which is attributed to the formation of some new grains with high angle boundaries. With an increase in the annealing time to 4 and 8 h, abnormal grain growth occurred, whereby the grain size reached up to several hundred micrometers and few millimeters (<xref ref-type="fig" rid="F6">Fig. 6</xref>(b, c)). Accordingly, the density of the boundary between 2°-15° dropped to zero, indicating no substructure within the grains, whereas a boundary density of greater than 15° decreased to close to zero (<xref ref-type="fig" rid="F7">Fig. 7</xref>(a)) owing to significant grain growth. ...</div> </div> </div> <div id="article_reference_meta_b32"> <div id="article_reference_meta_b32_title" class="title_"></div> <div id="article_reference_meta_b32_citedNumber">1</div> <div id="article_reference_meta_b32_nian"></div> <div id="article_reference_meta_b32_jcr"></div> <div id="article_reference_meta_b32_cjcr"></div> <div id="article_reference_meta_b32_articleCitedText"> <div class="sentence">... Regarding the three dispersoid-containing alloys, the number density of the dispersoids increased with an increase in the Mn content after homogenization (<xref ref-type="fig" rid="F3">Fig. 3</xref>). This increase in the number density leads to an increase in the substructure density after hot deformation (<xref ref-type="fig" rid="F4">Fig. 4</xref>, <xref ref-type="fig" rid="F5">Fig. 5</xref>). In theory, a high number density of dispersoids can contribute to a better recrystallization resistance during annealing owing to the strong pining ability on the grain boundary migration and the grain rotation [<xref ref-type="bibr" rid="b24">24</xref>,<xref ref-type="bibr" rid="b32">32</xref>]. However, a contrary result in the present study was observed: The 0.5 Mn alloy exhibited the best recrystallization resistance with the lowest SRX fraction, whereas the 0.75 Mn alloy possessed a higher SRX fraction and the 1 Mn alloy had even the highest (<xref ref-type="fig" rid="F6">Fig. 6</xref>, <xref ref-type="fig" rid="F8">Fig. 8</xref>(b)). ...</div> </div> </div> <div id="article_reference_meta_b33"> <div id="article_reference_meta_b33_title" class="title_"></div> <div id="article_reference_meta_b33_citedNumber">2</div> <div id="article_reference_meta_b33_nian">2010</div> <div id="article_reference_meta_b33_jcr"></div> <div id="article_reference_meta_b33_cjcr"></div> <div id="article_reference_meta_b33_articleCitedText"> <div class="sentence">... The decreased recrystallization resistance with an increase in the Mn content was believed to be related to the DFZs. <xref ref-type="fig" rid="F11">Fig. 11</xref> shows a bright-field TEM image of the recrystallized grains in the 1 Mn alloy after 8 h of annealing. Newly formed and recrystallized grains, featured as being free of internal substructures, can be clearly observed in the interdendrite region where large intermetallic particles are present. The location of these recrystallized grains was actually in the DFZ where nearly no dispersoids existed. However, in the neighbor regions, a large amount of dispersoids remained, representing a high density of substructures in the non-recrystallized grains. This result implies that during annealing, the newly recrystallized grains preferred to nucleate and grow at the DFZ where the pinning effect of dislocations is the weakest [<xref ref-type="bibr" rid="b33">33</xref>]. In addition, with the increase of the Mn content from 0.5 wt.% to 1 wt.%, the amount of intermetallic particles (Fe-rich intermetallic and primary Mg<sub>2</sub>Si) moderately increased [<xref ref-type="bibr" rid="b8">8</xref>] and those particles were primarily located in DFZ zones. The higher amount of large intermetallic particles in the higher Mn-containing alloys can also favor the recrystallization in the DFZ zones by the particle-stimulated nucleation mechanism in some extent [<xref ref-type="bibr" rid="b34">34</xref>]. Once the recrystallized grains encountered the dispersoid zone, the growth was arrested, and thus the growth of recrystallized grains was restricted in the DFZ. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... The coarsening and dissolution of the dispersoids occurred during the post-deformation' annealing in all three Mn-containing alloys owing to the less thermal stability of dispersoids at 500 °C (<xref ref-type="fig" rid="F9">Fig. 9</xref>, <xref ref-type="fig" rid="F10">Fig. 10</xref>). However, the above results indicate that the recrystallization resistance was mainly controlled by the DFZ fraction and less influenced by the number density and coarsening of the dispersoids. It was reported that even a low density of dispersoids of 0.003 μm<sup>-2</sup> can have a significant influence on the recrystallization resistance in Al-Mg-Si alloys [<xref ref-type="bibr" rid="b33">33</xref>]. The dispersoid densities after 8 h of annealing at 500 °C still ranged from 3.2-5.5 μm<sup>-2</sup> in the three Mn-containing alloys (<xref ref-type="fig" rid="F10">Fig. 10</xref>), which is probably sufficient for inhibiting recrystallization. Therefore, the distribution of dispersoids associated with the DFZ is in fact the predominant factor controlling the recrystallization resistance during the post-deformation annealing. ...</div> </div> </div> <div id="article_reference_meta_b34"> <div id="article_reference_meta_b34_title" class="title_"></div> <div id="article_reference_meta_b34_citedNumber">2</div> <div id="article_reference_meta_b34_nian"></div> <div id="article_reference_meta_b34_jcr"></div> <div id="article_reference_meta_b34_cjcr"></div> <div id="article_reference_meta_b34_articleCitedText"> <div class="sentence">... The decreased recrystallization resistance with an increase in the Mn content was believed to be related to the DFZs. <xref ref-type="fig" rid="F11">Fig. 11</xref> shows a bright-field TEM image of the recrystallized grains in the 1 Mn alloy after 8 h of annealing. Newly formed and recrystallized grains, featured as being free of internal substructures, can be clearly observed in the interdendrite region where large intermetallic particles are present. The location of these recrystallized grains was actually in the DFZ where nearly no dispersoids existed. However, in the neighbor regions, a large amount of dispersoids remained, representing a high density of substructures in the non-recrystallized grains. This result implies that during annealing, the newly recrystallized grains preferred to nucleate and grow at the DFZ where the pinning effect of dislocations is the weakest [<xref ref-type="bibr" rid="b33">33</xref>]. In addition, with the increase of the Mn content from 0.5 wt.% to 1 wt.%, the amount of intermetallic particles (Fe-rich intermetallic and primary Mg<sub>2</sub>Si) moderately increased [<xref ref-type="bibr" rid="b8">8</xref>] and those particles were primarily located in DFZ zones. The higher amount of large intermetallic particles in the higher Mn-containing alloys can also favor the recrystallization in the DFZ zones by the particle-stimulated nucleation mechanism in some extent [<xref ref-type="bibr" rid="b34">34</xref>]. Once the recrystallized grains encountered the dispersoid zone, the growth was arrested, and thus the growth of recrystallized grains was restricted in the DFZ. ...</div> <div class="boundary"><p class="ty-x"></p></div> <div class="sentence">... <xref ref-type="fig" rid="F13">Fig. 13</xref> shows a schematic of how the recrystallization took place. After a hot deformation, numerous substructures were induced at both the dispersoids zones and DFZs (<xref ref-type="fig" rid="F13">Fig. 13</xref>(a)). During post-deformation annealing, the high temperature provided a driving force for the motion of dislocations and subgrains. In the dispersoid zone, owing to the strong pinning effect of the dispersoids on the dislocations and subgrain boundaries, only SRV was able to take place. However, in the DFZs, SRX can start through a diminishment or coalescence of dislocations into the subgrains owing to the absence of dispersoids and the weakest pinning effect (<xref ref-type="fig" rid="F13">Fig. 13</xref>(b)). In addition, within the PFZs, the region surrounding the large intermetallic particles was highly strained during hot deformation, resulting in a higher density of dislocations and subgrains compared to the other regions. During annealing, the nucleation of new grains preferred to occur near the intermetallic particles where the driving force was higher, which is known as a particle-stimulated nucleation of recrystallization [<xref ref-type="bibr" rid="b2">2</xref>,<xref ref-type="bibr" rid="b34">34</xref>,<xref ref-type="bibr" rid="b35">35</xref>]. As a result, DFZs acted as the preferred regions where the SRX started and propagated (<xref ref-type="fig" rid="F13">Fig. 13</xref>(c)). ...</div> </div> </div> <div id="article_reference_meta_b35"> <div id="article_reference_meta_b35_title" class="title_"></div> <div id="article_reference_meta_b35_citedNumber">1</div> <div id="article_reference_meta_b35_nian"></div> <div id="article_reference_meta_b35_jcr"></div> <div id="article_reference_meta_b35_cjcr"></div> <div id="article_reference_meta_b35_articleCitedText"> <div class="sentence">... <xref ref-type="fig" rid="F13">Fig. 13</xref> shows a schematic of how the recrystallization took place. After a hot deformation, numerous substructures were induced at both the dispersoids zones and DFZs (<xref ref-type="fig" rid="F13">Fig. 13</xref>(a)). During post-deformation annealing, the high temperature provided a driving force for the motion of dislocations and subgrains. In the dispersoid zone, owing to the strong pinning effect of the dispersoids on the dislocations and subgrain boundaries, only SRV was able to take place. However, in the DFZs, SRX can start through a diminishment or coalescence of dislocations into the subgrains owing to the absence of dispersoids and the weakest pinning effect (<xref ref-type="fig" rid="F13">Fig. 13</xref>(b)). In addition, within the PFZs, the region surrounding the large intermetallic particles was highly strained during hot deformation, resulting in a higher density of dislocations and subgrains compared to the other regions. During annealing, the nucleation of new grains preferred to occur near the intermetallic particles where the driving force was higher, which is known as a particle-stimulated nucleation of recrystallization [<xref ref-type="bibr" rid="b2">2</xref>,<xref ref-type="bibr" rid="b34">34</xref>,<xref ref-type="bibr" rid="b35">35</xref>]. As a result, DFZs acted as the preferred regions where the SRX started and propagated (<xref ref-type="fig" rid="F13">Fig. 13</xref>(c)). ...</div> </div> </div> <div id="article_reference_meta_b36"> <div id="article_reference_meta_b36_title" class="title_"></div> <div id="article_reference_meta_b36_citedNumber">1</div> <div id="article_reference_meta_b36_nian"></div> <div id="article_reference_meta_b36_jcr"></div> <div id="article_reference_meta_b36_cjcr"></div> <div id="article_reference_meta_b36_articleCitedText"> <div class="sentence">... It is worthwhile to note that for structural applications, the hot-deformed 6082 alloys after post-deformation annealing/solution usually undergo an artificial aging to precipitate the nanoscale β”/β'-Mg<sub>2</sub>Si phases and to achieve the adequately high strengths. The presence of a large number of dispersoids induced by Mn addition can consume a part of Si solutes, which may disfavor β”/β' precipitation. On the other side, the pre-existing dispersoids formed during homogenization could provide favorable nucleation sites for subsequent β”/β' precipitation owing to the multiple benefits of the nucleation effects between the dispersoids and Mg<sub>2</sub>Si [<xref ref-type="bibr" rid="b36">36</xref>]. In addition to enhanced recrystallization resistance, it would be expected that the presence of dispersoids further improves the mechanical properties of final products. ...</div> </div> </div> </div> </div> </div> </div> <div class="cankaowenxian1"></div> <div class="col-xs-3"> <div class="slide" style="top: 367px; display: block;"> <div id="sideToolbar"> <div id="sideCatalog" class="sideCatalogBg"> <div id="sideCatalog-sidebar"> <div class="sideCatalog-sidebar-top"></div> <div class="sideCatalog-sidebar-bottom"></div> </div> <div id="sideCatalog-updown"> <div id="sideCatalog-up" class="sideCatalog-up-disable" onclick="shang(this)" title="up"></div> <div id="sideCatalog-down" class="sideCatalog-down-enable" title="down"></div> </div> <div id="sideCatalog-catalog"> <dl style="width: 175px; zoom: 1; top: 0px;" id="outline_dl"> </dl> </div> </div> <i id="sideToolbar-up" style="float:left;"></i> <i id="sideCatalogBtn" style="float:left;"></i> </div> </div> </div> </div> </div> </div> </div> <div class="motaikuang" onclick="guanbiann()"></div> <div class="tupiankuang"> <i class="glyphicon glyphicon-remove-circle guanbi" onclick="guanbiann();"></i> <p class="biaotitupiana"> <span class="yige"></span>/<span class="zongdea"></span> </p> <div class="imgaesa"> <table class="haoshi"> <tr> <td onclick="shang();" class="zuo" > 〈 </td> <td class="zhong"> <img src="" class="tp" onload="xianss2()"> </td> <td onclick="xia();" class="you" > 〉 </td> </tr> </table> </div> <div class="neirong"></div> </div> <div class="biaokuang"> <i class="glyphicon glyphicon-remove-circle guanbi" onclick="guanbiann();"></i> <div class="tuti"></div> <div class="table-responsive liuliu"></div> <div class="biaoshuo"> <p class="mtk-biaoshoen"></p> <p class="mtk-biaoshocn"></p> </div> </div> <div class="qipao-zong" onmouseleave="yichu()"> <div class="qipao-content"> </div> <div class="qipao-jiantou"> <img src="../../../richhtml/1005-0302/richHtml_jats1_1_en/images/1.png"> </div> </div> </div> </div> <div class="container"> <!--footer--> <div class="row footer-box"> <dl class="col-md-3"> <dt>About JMST</dt> <li><a href="https://www.jmst.org/EN/column/column148.shtml">About JMST</a></li> <li><a href="https://www.jmst.org/EN/column/column149.shtml">Editorial board</a></li> <li><a href="https://www.jmst.org/EN/column/column260.shtml">Coverage</a></li> <li><a href="https://www.jmst.org/EN/column/column261.shtml">Index in</a></li> </dl> <dl class="col-md-3"> <dt>Authors and reviewers</dt> <li><a href="https://www.jmst.org/EN/column/column150.shtml">How to submit</a></li> <li><a href="https://www.jmst.org/EN/column/column152.shtml">Art instruction</a></li> <li><a href="https://www.editorialmanager.com/j-mst" target="_blank">Submission online</a></li> <li><a href="https://www.jmst.org/EN/column/column265.shtml">Instruction for reviewer</a></li> <li><a href="https://www.editorialmanager.com/j-mst" target="_blank">Reviewer login</a></li> <li><a href="https://www.editorialmanager.com/j-mst" target="_blank">Office login</a></li> </dl> <dl class="col-md-4"> <dt>Contact</dt> <li>Tel: +86-024-83978208<br> E-mail: jmst@imr.ac.cn<br> Address: 72 Wenhua Road, Shenyang, Liaoning, 110016, China </li> </dl> <dl class="col-md-2 text-center"> <img src="https://www.jmst.org/images/1005-0302/images/ewm.jpg" class="img-responsive" style="width:104px;border-radius: 1px;margin:0 auto"> WeChat </dl> <div class="col-md-12"> <ul style="line-height:160%;text-align:center;font-size:12px;"> <li>Copyright © Editorial Office of Journal of Materials Sciences and Technology, All Rights Reserved.</li> <script type="text/javascript">document.write(unescape("%3Cspan id='_ideConac' %3E%3C/span%3E%3Cscript src='http://dcs.conac.cn/js/33/000/0000/60427449/CA330000000604274490014.js' type='text/javascript'%3E%3C/script%3E"));</script> </ul> </div> </div> </div> <!--返回顶部--> <div class="top_web" id="backtop" style="display:block;"> <span class="glyphicon glyphicon-chevron-up" aria-hidden="true" ></span> </div> <script src="https://www.jmst.org/images/1005-0302/js/backtop.js"></script> <!--返回顶部end--> </body> <script type="text/javascript" src="../../../richhtml/1005-0302/richHtml_jats1_1_en/js/wangkan.js"></script> <script src="../../../richhtml/1005-0302/richHtml_jats1_1_en/js/qrcode.js"></script> <script type="text/x-mathjax-config"> $(function() { $(".keyword").each(function() { // key var key = $(this); // value var value = key.text(); // 添加点击事件 $(key).click(function() { // 创建一个form表单 var form = document.createElement("form"); form.method = "post"; form.action = "https://www.jmst.org/EN/article/advancedSearchResult.do"; form.target = "_blank"; // 将form表单添加到body标签中 document.body.appendChild(form); // 新建表单标签 var formElement = document.createElement("input"); formElement.name = "searchSQL"; formElement.type = "hidden"; formElement.value = "(" + value + "[Keyword])"; form.appendChild(formElement); form.submit(); }); }); }); </script> </html>