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(a) Microstructures of the AlÀAl 2 Cu eutectic, the dark phase is a-Al and the light phase is u-Al 2 Cu. (b) Localized shears in rolled nanoscale Al 2 Cu lamellae of the laser surface re-melting treated material. (c) Diffraction patterns of the Al 2 Cu lamella from "B" area in (b), showing the shear plane is (011)Al 2 Cu. (d) The interface plane associated with variant I OR and (e) the corresponding diffraction patterns of the two phases, indicating the interface plane is (001)Al 2 Cu||(001)Al. (f) HRTEM image showing a terraced (001)Al 2 Cu||(001)Al interface structure.
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Laser surface re-melted Al–Al2Cu eutectic alloy with α-Al and θ-Al2Cu nanoscale lamellae exhibits high strength and good plasticity at room temperature, implying that the nanoscale θ-Al2Cu lamellae plastically co-deform with α-Al. Microscopy characterization reveal that plastic deformation of θ-Al2Cu lamellae is accommodated by localized shear on u...
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... Al-Cu eutectic alloys exhibit microscale eutectics comprising alternate Al and Al 2 Cu lamina with the »1À3 mm interlamellar spacing. Through laser surface re-melting AlÀAl 2 Cu eutectics, the interlamellar spacing can be refined to »40 nm (Fig. 1a) [9]. Nanoscale Al-A 2 Cu eutectic shows the simultaneous improvement in their strength and ductility [18,40]. Plastic co-deformation between Al 2 Cu and Al layers was observed in rolled samples by transmission electron microscopy [18]. The nanoscale Al 2 Cu lamellae plastically deform via localized shear on {011} slip plane that are ...
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... shows the simultaneous improvement in their strength and ductility [18,40]. Plastic co-deformation between Al 2 Cu and Al layers was observed in rolled samples by transmission electron microscopy [18]. The nanoscale Al 2 Cu lamellae plastically deform via localized shear on {011} slip plane that are not observed in monolithic crystals of Al 2 Cu (Figs. ...
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... plane could be either (001) Al2Cu ||{001} Al or {121} Al2Cu ||{111} Al . For variant II, the interface plane is {210} Al2Cu ||{110} Al . In nanoscale AlÀAl 2 Cu eutectics, we found that variant I prevailed over variant II [9], and (001) Al2Cu || {001} Al interface plane has the lowest interface energy of 529 mJ/m 2 among other interface planes. Fig. 1e (001) Al terrace and steps (Fig. 1f). Therefore, the corresponding OR is described with the interface plane (001) [310] Al directions (denoted by the x'-and z'-axis) respectively, as shown in Fig. 2a. Considering the <110> Al direction is the close packed direction in fcc-Al, the <210> Al2Cu ||<110> Al axes will be used as reference ...
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... ||{001} Al or {121} Al2Cu ||{111} Al . For variant II, the interface plane is {210} Al2Cu ||{110} Al . In nanoscale AlÀAl 2 Cu eutectics, we found that variant I prevailed over variant II [9], and (001) Al2Cu || {001} Al interface plane has the lowest interface energy of 529 mJ/m 2 among other interface planes. Fig. 1e (001) Al terrace and steps (Fig. 1f). Therefore, the corresponding OR is described with the interface plane (001) [310] Al directions (denoted by the x'-and z'-axis) respectively, as shown in Fig. 2a. Considering the <110> Al direction is the close packed direction in fcc-Al, the <210> Al2Cu ||<110> Al axes will be used as reference coordinate axes in the following ...
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... initiation of local shears in Al 2 Cu layers. Under this loading condition, the eight {111}<110> slip systems in fcc-Al crystal could be equally activated because of the same Schmid factor of 0.4083 (Fig. 8). MD simulations show that plastic deformation commences in Al layers ( Fig. 9) and then local shears on {211} Al2Cu planes in Al 2 Cu layers (Fig. ...
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... configurations corresponding to A-E in the stress-strain curve are shown in Fig. 9 and Fig. 10, where the atoms are colored by the local von Mises shear strain invariant [64], which is calculated from the local deformation gradient with respect to the zero applied compressive strain. At small compressive strain below a critical strain about 7.6%, the stress almost linearly increases with the applied strain, corresponding to the ...
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... deformation. With continuous normal compression after exceeding the critical strain of 7.6%, the stress-strain curve deviates from elastic response, which is associated with plastic deformation in the Al layer by nucleating a/2<110> dislocation loops from the segments of interface dislocations, as can be seen in Fig. 9b and the illustration in Fig. 9b 1 and 9b 2 . The nucleated dislocations glide on the four symmetrical {111} planes (red solid curves) in Al layers, depositing dislocations (blue solid lines) on the interface. As dislocations glide on the four symmetrical {111} planes in the Al layer, some dislocation segments could be annihilated on the intersections of two different {111} ...
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... is noted that these dislocations at the two interfaces have the same Burgers vector but the opposite line sense (blue solid and red solid lines on the interfaces shown in Fig. 9b 1 -9b 3 ). A strong attractive force is thus generated and acted on each other. More importantly, a maximum shear stress will be on planes in Al 2 Cu layer which are nearly parallel to {111} planes in Al layers [13]. Once the resolved shear stress (contribution from both the applied stress and the stress field of accumulated dislocations on the ...
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... Once the resolved shear stress (contribution from both the applied stress and the stress field of accumulated dislocations on the opposite sides of the interface) exceeds the critical resolved shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al ...
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... exceeds the critical resolved shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with ...
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... shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> ...
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... is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to ...
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... the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer ...
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... (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results ...
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... in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results confirm that the slip occurring in Al 2 Cu layers is due to the continuity of slip systems across the ...
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... 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results confirm that the slip occurring in Al 2 Cu layers is due to the continuity of slip systems across the ...
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... metal and intermetallic layers, which is the origin of high strain hardening rate on the stress-strain curve in Fig. 7b. With accumulation of dislocations in the interface plane, the local shear stresses associated with these accumulated dislocations and the applied stress will activate glide dislocations in Al 2 Cu layer [13], as illustrated in Fig. 10a 2 . As a result, unusual shear associated with {121}<111> slip system takes place in nanoscale Al 2 Cu layer ( Fig. 10a and ...
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... With accumulation of dislocations in the interface plane, the local shear stresses associated with these accumulated dislocations and the applied stress will activate glide dislocations in Al 2 Cu layer [13], as illustrated in Fig. 10a 2 . As a result, unusual shear associated with {121}<111> slip system takes place in nanoscale Al 2 Cu layer ( Fig. 10a and ...
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... Al-Cu eutectic alloys exhibit microscale eutectics comprising alternate Al and Al 2 Cu lamina with the »1À3 mm interlamellar spacing. Through laser surface re-melting AlÀAl 2 Cu eutectics, the interlamellar spacing can be refined to »40 nm (Fig. 1a) [9]. Nanoscale Al-A 2 Cu eutectic shows the simultaneous improvement in their strength and ductility [18,40]. Plastic co-deformation between Al 2 Cu and Al layers was observed in rolled samples by transmission electron microscopy [18]. The nanoscale Al 2 Cu lamellae plastically deform via localized shear on {011} slip plane that are ...
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... shows the simultaneous improvement in their strength and ductility [18,40]. Plastic co-deformation between Al 2 Cu and Al layers was observed in rolled samples by transmission electron microscopy [18]. The nanoscale Al 2 Cu lamellae plastically deform via localized shear on {011} slip plane that are not observed in monolithic crystals of Al 2 Cu (Figs. ...
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... plane could be either (001) Al2Cu ||{001} Al or {121} Al2Cu ||{111} Al . For variant II, the interface plane is {210} Al2Cu ||{110} Al . In nanoscale AlÀAl 2 Cu eutectics, we found that variant I prevailed over variant II [9], and (001) Al2Cu || {001} Al interface plane has the lowest interface energy of 529 mJ/m 2 among other interface planes. Fig. 1e (001) Al terrace and steps (Fig. 1f). Therefore, the corresponding OR is described with the interface plane (001) [310] Al directions (denoted by the x'-and z'-axis) respectively, as shown in Fig. 2a. Considering the <110> Al direction is the close packed direction in fcc-Al, the <210> Al2Cu ||<110> Al axes will be used as reference ...
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... ||{001} Al or {121} Al2Cu ||{111} Al . For variant II, the interface plane is {210} Al2Cu ||{110} Al . In nanoscale AlÀAl 2 Cu eutectics, we found that variant I prevailed over variant II [9], and (001) Al2Cu || {001} Al interface plane has the lowest interface energy of 529 mJ/m 2 among other interface planes. Fig. 1e (001) Al terrace and steps (Fig. 1f). Therefore, the corresponding OR is described with the interface plane (001) [310] Al directions (denoted by the x'-and z'-axis) respectively, as shown in Fig. 2a. Considering the <110> Al direction is the close packed direction in fcc-Al, the <210> Al2Cu ||<110> Al axes will be used as reference coordinate axes in the following ...
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... initiation of local shears in Al 2 Cu layers. Under this loading condition, the eight {111}<110> slip systems in fcc-Al crystal could be equally activated because of the same Schmid factor of 0.4083 (Fig. 8). MD simulations show that plastic deformation commences in Al layers ( Fig. 9) and then local shears on {211} Al2Cu planes in Al 2 Cu layers (Fig. ...
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... configurations corresponding to A-E in the stress-strain curve are shown in Fig. 9 and Fig. 10, where the atoms are colored by the local von Mises shear strain invariant [64], which is calculated from the local deformation gradient with respect to the zero applied compressive strain. At small compressive strain below a critical strain about 7.6%, the stress almost linearly increases with the applied strain, corresponding to the ...
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... deformation. With continuous normal compression after exceeding the critical strain of 7.6%, the stress-strain curve deviates from elastic response, which is associated with plastic deformation in the Al layer by nucleating a/2<110> dislocation loops from the segments of interface dislocations, as can be seen in Fig. 9b and the illustration in Fig. 9b 1 and 9b 2 . The nucleated dislocations glide on the four symmetrical {111} planes (red solid curves) in Al layers, depositing dislocations (blue solid lines) on the interface. As dislocations glide on the four symmetrical {111} planes in the Al layer, some dislocation segments could be annihilated on the intersections of two different {111} ...
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... is noted that these dislocations at the two interfaces have the same Burgers vector but the opposite line sense (blue solid and red solid lines on the interfaces shown in Fig. 9b 1 -9b 3 ). A strong attractive force is thus generated and acted on each other. More importantly, a maximum shear stress will be on planes in Al 2 Cu layer which are nearly parallel to {111} planes in Al layers [13]. Once the resolved shear stress (contribution from both the applied stress and the stress field of accumulated dislocations on the ...
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... Once the resolved shear stress (contribution from both the applied stress and the stress field of accumulated dislocations on the opposite sides of the interface) exceeds the critical resolved shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al ...
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... exceeds the critical resolved shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with ...
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... shear stress associated with slip in Al 2 Cu layer, localized shears in Al 2 Cu layers occur, as indicated by the localized Mises shear strain shown in Fig. 10a. The localized shears of (211) Al2Cu (which is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> ...
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... is the 1st slip pair in Fig. 8b). The two dislocations meet in the middle of Al 2 Cu layer, as magnified in Fig. 10a 1 and illustrated in Fig. 10a 2 . Symmetrically, there are another two {121}<111> slip systems in Al 2 Cu that are coupled with the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to ...
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... the dislocation slips on (111) Al and (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer ...
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... (111) Al planes respectively, as indicated by the white circle in Fig. 10b and the 2nd and 4th slip pairs illustrated in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results ...
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... in Fig. 8. The nucleated dislocations in Al 2 Cu with Burgers vector of a/2<111> Al2Cu on {211} Al2Cu planes can also propagate through the Al 2 Cu layer, as indicated by the white circles in Fig. 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results confirm that the slip occurring in Al 2 Cu layers is due to the continuity of slip systems across the ...
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... 10c. The magnification shown in Fig. 10c 2 clearly displays the local shears of Al 2 Cu due to the accumulated dislocations in the interface (Fig. 10a 2 ). This process will leave slip traces on the interface, which are parallel to [110] Al2Cu , as denoted by the traces circled in the black lines on the atomic structures of Al 2 Cu layer shown in Fig. 10c 1 . These results confirm that the slip occurring in Al 2 Cu layers is due to the continuity of slip systems across the ...
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... metal and intermetallic layers, which is the origin of high strain hardening rate on the stress-strain curve in Fig. 7b. With accumulation of dislocations in the interface plane, the local shear stresses associated with these accumulated dislocations and the applied stress will activate glide dislocations in Al 2 Cu layer [13], as illustrated in Fig. 10a 2 . As a result, unusual shear associated with {121}<111> slip system takes place in nanoscale Al 2 Cu layer ( Fig. 10a and ...
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... With accumulation of dislocations in the interface plane, the local shear stresses associated with these accumulated dislocations and the applied stress will activate glide dislocations in Al 2 Cu layer [13], as illustrated in Fig. 10a 2 . As a result, unusual shear associated with {121}<111> slip system takes place in nanoscale Al 2 Cu layer ( Fig. 10a and ...
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... 2a and 9). Deformability and deformation mechanisms for nanoscale sandwiched intermetallic branches could differ from their bulk counterpart, due to the discrepancy in scale and confined loading state, as corroborated by studies on laminated nanolayers [65][66][67][68] . Comparing to transient and reversible GND pileups adjacent to Alintermetallic interfaces, temporary defects configuration existing in intermetallic-intermetallic interface could generate back stress as well. ...
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... In the introduction, we highlighted the structural similarity between the Ω unit cell and the θ(Al2Cu) when the cell axis is redefined [10]. We therefore summarized the reported slip systems for the θ phases in the Al-Al 2 Cu eutectic alloy [51,55] and micropillar compression on the θ intermetallic [45] in Fig. 7 [51], might show a lower energy barrier compared to our GSFE results, it is not activated in Ω precipitates because of the low geometric compatibility. Furthermore, none of the slip systems activated in the micropillar compression study [45] were observed in our study. ...
... Consider a single NP in a GB that is modeled by a long dilatational inclusion with cross section of rectangular shape ABCD, the dimensions of which are given by the diagonal h and the angle α (Figure 1a). It is well known [28][29][30][31][32] that the larger faces of Al 2 Cu NPs in aluminum-based alloys lie along {111} planes of the aluminum matrix. This can be explained by the growth kinetics of the Al 2 Cu intermetallic compound, whose {110} faces grow faster than the others [28], and by a relatively low lattice misfit f at such boundary [32]. ...
... It is well known [28][29][30][31][32] that the larger faces of Al 2 Cu NPs in aluminum-based alloys lie along {111} planes of the aluminum matrix. This can be explained by the growth kinetics of the Al 2 Cu intermetallic compound, whose {110} faces grow faster than the others [28], and by a relatively low lattice misfit f at such boundary [32]. It is assumed that the boundaries of the NP are initially in a coherent state, that is, they do not contain misfit dislocations. ...
... Consider a single NP in a GB that is modeled by a long dilatational inclusion with cross section of rectangular shape ABCD, the dimensions of which are given by the diagonal h and the angle α (Figure 1a). It is well known [28][29][30][31][32] that the larger faces of Al2Cu NPs in aluminum-based alloys lie along {111} planes of the aluminum matrix. This can be explained by the growth kinetics of the Al2Cu intermetallic compound, whose {110} faces grow faster than the others [28], and by a relatively low la>ice misfit f at such boundary [32]. ...
A theoretical model is suggested that describes the effect of deformation temperature on the yield stress of an ultrafine-grained (UFG) Al-Cu-Zr alloy structured with severe plastic deformation. Within the model, nanoprecipitates (NPs) of Al2Cu act as sources of lattice dislocations in the presence of a number of extrinsic grain-boundary dislocations (EGBDs) near the NPs. It is shown that the number of EGBDs near the NPs decreases with a drop in the deformation temperature that increases the yield stress of the Al-Cu-Zr alloy. The proposed model is in good quantitative agreement with available experimental results.
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... Following the authors of Ref. [31], consider a separate NP in a GB, modeled by the inclusion of a rectangular shape ABCD, the sizes of which are given by the diagonal h and the angle of inclination α of the BC face to the GB plane (Fig. 20). It is known [73][74][75][76][77] that the formation of lamellar Al2Cu-NPs in Al-based alloys 1.71%) f = directions, as well as by the growth kinetics of the Al2Cu intermetallic, whose {110} faces grow faster than the others [73]. Note that recent works on computer simulation of the structure and properties of such interfaces in lamellar eutectic Al-Al2Cu composites [75,76] have shown that they relatively easily transit from a coherent state to a semicoherent one due to the filling of the interfaces with three families of misfit dislocations (MDs) by sliding along the interface of partial Shockley dislocations. ...
... Based on the results of experimental observations and computer simulations of the Al2Cu NPs in aluminum alloys [73][74][75][76][77], the authors of Ref. [31] assumed that the easy slip plane of LDs coincided with the plane of the BC face and made an angle α with the GB plane. The emission of such a b-dislocation was represented as the nucleation of an LD dipole with Burgers vectors ±b (Fig. 20). ...
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This is a brief review of recent experimental and theoretical results on the influence of low temperature annealing and subsequent small plastic deformation on microstructure, strength and ductility of ultrafine-grained Al and Al-based alloys structured by high pressure torsion. Some earlier results on this problem for ultrafine-grained Al and Al-based alloys structured by different methods of severe plastic deformation are also shortly presented. The reasons for the effects of hardening by annealing and softening by additional small plastic deformation of the materials are suggested and discussed in detail. Moreover, the influence of the temperature of mechanical testing and the alloying elements are in the focus of the review. It is shown that in the physical origin of these effects are the transformations of the defect structure of grain boundaries in the process of low temperature annealing and subsequent small plastic deformation of the ultrafine-grained Al and Al-based alloys structured by high pressure torsion.
... Observation de grains en ilôts ou disques en (a) imagerie électronique et par (b) EBSD sur la condition P250V1250e30 90h120.A fort grossissement, de fines lamelles blanches micrométriques sont révélées à l'intérieur des grains (Figure 3.18a et 3.18b). Cette structure lamellaire rappelle les structures eutectiques classiques, tels que l'Al/Al2Cu[Guisen+2020]. Cette microstructure suggèrerait une ségrégation des éléments Ni et Ti pendant la solidification. ...
Depuis les années 90, les outils endodontiques utilisés par les dentistes ont évolué de l’acier vers le NiTi pour exploiter les propriétés superélastiques de cet alliage à mémoire de forme. Ces outils très fins sont rapidement endommagés et mis hors d'usage par des sollicitations cycliques complexes (fatigue en traction/compression, torsion, usure, etc.). Actuellement, ces outils sont fabriqués par usinage de fils tréfilés (diamètre 0.1 à 2 mm) provoquant un grand nombre de rebuts et limitant la liberté sur les géométries fabriquées. Le procédé de fusion laser (SLM) sur lit de poudre métallique est un procédé avancé permettant de réaliser des géométries complexes aux dimensions micrométriques et pourrait convenir à la fabrication de tels outils, en apportant plus de liberté dans leur conception. Cette thèse s’attache à comprendre les relations entre les paramètres du procédé SLM, la microstructure et les propriétés mécaniques du NiTi-SLM à l'état brut de fabrication, afin d’identifier des paramètres de fabrication optimaux. Un intervalle de fusion dit "acceptable", décrit par des termes d'énergies volumiques, a été identifié à l'aide d'analyses détaillées de la taille des grains, des profondeurs des micro-bains de fusion et du taux de porosité dans le matériau. Une approche analytique considérant la conduction thermique de la chaleur a été utilisée pour estimer la profondeur de ces bains et prédire les paramètres optimaux de fusion. L’influence des paramètres de fabrication sur les phases, les transformations de phases et les précipitations présentes dans le matériau a aussi été étudiée. Elle a mis en évidence 6 structures cristallographiques principales : NiTiB2, Ni3Ti > Ni4Ti3, NiTiB19’, Ni3TiO5 et NiTi2. L'élargissement important des pics en DRX et DSC a été expliqué par la présence d'oxydes associés à ces 6 phases, à la génération de distorsions/contraintes thermomécaniques lors de la fabrication et à une distribution de composition chimique sur l'ensemble du matériau. La vitesse de balayage laser a révélé qu'elle génère une forte distorsion sur les paramètres de maille. Enfin, les propriétés mécaniques en traction cyclique des échantillons NiTi-SLM se sont révélées très inférieures à celles des NiTi-tréfilés. En effet, la microstructure NiTi-SLM engendre un faible taux d'amortissement et un fort endommagement plastique par des microbandes de cisaillement à travers l'échantillon et menant à une rupture mixte fragile-ductile.
... Further studies on the activated plasticity mechanisms of the pure intermetallic phase are mostly conducted at elevated temperatures, whereas studies on the lamellar Al-Al2Cu eutectic allow to plastically deform the compound at ambient temperature. These studies revealed dislocations in the [3,4,[11][12][13][14] as well as localized shear on {011} planes and faulted structures on {211} planes [12]. A recent publication by the present authors [15] further took advantage of nanomechanical methods to access the plasticity of the Al2Cu θ-phase at ambient temperature. ...
The investigation of the deformation behaviour of intermetallic phases is mostly limited to high temperatures due to their low ductility at ambient temperature. Therefore, within this study, nanoindentation experiments on the Al$_{2}$Cu phase were performed from ambient temperature up to 300{\deg}C in conjunction with TEM investigations of the deformed material. It was found that the Al$_{2}$Cu phase starts to soften from temperatures above 150{\deg}C. The low temperature deformation behaviour was dominated by serrated yielding and transitioned towards a smooth deformation behaviour above 150{\deg}C. No slip traces were observed in the vicinity of the indents in both temperature regimes. The predominant dislocations observed after ambient temperature deformation and elevated temperature deformation are assumed to belong to the same operating slip systems. Therefore, the underlying deformation mechanism for both temperatures appears to be the same and thermally activated.
... Recently, Choudhuri et al. [36,37] used MD simulations to study the effect of interfaces of the Ni/NiAl intermetallic composites on uniaxial deformation mechanisms, which depends on the B2-NiAl layer thickness. In another MD simulation study, it is found that in nanoscale composites Al-Al 2 Cu, plasticity is accommodated by localized shear on unusual slip planes, which is because of the continuity of slip systems across the interfaces [38]. A further systematic atomistic modeling study on the nanolayered composites comprised of intermetallic interfaces and their deformation responses will help to better understand the mechanical response of such material systems. ...
In this work, the deformation response of the B2-FeAl/Al intermetallic composites, as a model material system for nanolayered composites comprised of intermetallic interfaces, has been explored. We use atomistic simulations to study the deformation mechanisms and the interface misfit dislocation structure of B2-FeAl/Al nanolayered composites. It is shown that two sets of dislocations are contained in the interface misfit dislocation network and are correlated with the initial dislocation nucleation from the interfaces. The effects of layer thickness on the uniaxial deformation response of the B2-FeAl/Al multilayers are investigated. We observed that under compressive loading the smaller proportion of the FeAl layers leads to the lower overall flow stress. Under tensile loading, the void formation mechanism is investigated, suggesting the interface structure and the dislocation activities in the FeAl layers playing a significant role to trigger the strain localization which leads to void nucleation commencing at the interface. It is also found that the deformation behavior in the “weak” Fe/Cu interface behaves substantially different than that of the “strong” FeAl/Al interface. The atomistic modeling study of the nanolayered composites here underpinned the mechanical response of “strong” intermetallic interface material systems. There is no void nucleation during the entire plastic deformations in the Fe/Cu simulations, which is attributed to much higher dislocation density, more slip systems activated, and relative uniformly distributed dislocation traces in the Fe phase of the Fe/Cu multilayers.
... This was partially attributed to the refined θ-Al 2 Cu network in the AM alloy, whereas the brittle and coarse θ-Al 2 Cu precipitates along grain boundaries in the cast alloy acted as crack initiation sites that limited the tensile ductility [299]. New deformation mechanisms involving co-deformation of the matrix and nanoscale θ-Al 2 Cu lamellae are reported to contribute to excellent plasticity in laser re-melted Al-Cu alloys of eutectic composition (32.7 wt-%) [300][301][302]. Although the θ-Al 2 Cu is expected to coarsen significantly at elevated temperatures given the high diffusivity of Cu, these results offer exciting promise for increased ductility in AM HiFI alloys with nanoscale intermetallic networks. ...
Research on powder-based additive manufacturing of aluminium alloys is rapidly increasing, and recent breakthroughs in printing of defect-free parts promise substantial movement beyond traditional Al–Si–Mg) systems. One potential technological advantage of aluminium additive manufacturing, however, has received little attention: the design of alloys for use at T > ~200°C, or ~1/2 of the absolute melting temperature of aluminium. Besides offering lightweighting and improved energy efficiency through replacement of ferrous, titanium, and nickel-based alloys at 200–450°C, development of such alloys will reduce economic roadblocks for widespread implementation of aluminium additive manufacturing. We herein review the existing additive manufacturing literature for three categories of potential high-temperature alloys, discuss strategies for optimizing microstructures for elevated-temperature performance, and highlight gaps in current research. Although extensive microstructural characterisation has been performed on these alloys, we conclude that evaluations of their high-temperature mechanical properties and corrosion responses are severely deficient.
