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Microstructure evolution and deformation features of AZ31 Mg-alloy during creep
Materials Science and Engineering A 415 (2006) 309–316
Microstructure evolution and deformation features of AZ31 Mg-alloy during creep Sugui Tian a,∗ , Ling Wang a , Keun Yong Sohn b , Kyung Hyun Kim b , Yongbo Xu c , Zhuangqi Hu c a
School of Materials Science and Engineering, Shenyang University of Technology, No. 58 Xinghua South Street Tiexi District, Shenyang 110023, China b Korea Institute of Machinery and Materials, Changwon 641-010, South Korea c State Key Laboratory for Fatigue and Fracture of Materials, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China Received in revised form 3 October 2005; accepted 3 October 2005
Abstract By means of the measurement of the creep curve and the observation of SEM and transmission electron microscope (TEM), an investigation has been made into the microstructure evolution and deformation features of AZ31 Mg-alloy during high temperature creep. Results show that the deformation features of the alloy in the primary stage of creep are that significant amount of dislocation slips are activated on basal and non-basal planes, then these ones are concentrated into the dislocation cells or walls as creep goes on. At the same time, twinning occurs as an additional deformation mechanism in the role of the compatibility stress. During steady state creep, the dislocation cells are transformed into the subgrains, then, the protrusion and coalition of the sub-boundaries results in the occurrence of dynamic recovery (DRV). After the dynamic recrystallization (DRX), the multiple slips in the grain interiors are considered to be the main deformed mechanism in the later stage of the steady state creep. An obvious feature of creep entering the tertiary stage is that the cracks appear on the locations of the triple junction. As creep continues, the cracks are viscous expanded along the grain boundaries; this is taken for being the fracture mechanism of the alloy crept to failure. The multiple slips in the grain interiors and the cracks expanded viscous along the grain boundary occur in whole of specimens, that, together with the twins and dynamic recrystallization, is responsible for the rapid increase of the strain rate in the later stage during creep. © 2005 Elsevier B.V. All rights reserved. Keywords: AZ31 Mg-alloy; Creep; Dislocation slip; Twinning; Dynamic recrystallization; Deformation mechanism
1. Introduction Magnesium alloys are emerging as potentially good candidates for numerous applications, especially in the automotive industry. Their good properties, such as low density and high specific strength, make them promising replacements for other heavier materials like, for instance, steel, cast iron, and even aluminium [1]. However, the disadvantage of Mg-alloys is that they exhibit only limited ductility due to their HCP structure. Therefore, significant development effort is needed in order to widen the application of these materials. Magnesium alloys are being used to an increasing extent in applications where the components are subjected to elevated temperatures. Consequently, research is being focused on the
∗
Corresponding author. Tel.: +86 24 25694659; fax: +86 24 25691802. E-mail address: [email protected] (S. Tian).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.10.015
development of alloys able to withstand high stresses at temperature ranges of up to 180, 200 and 250 ◦ C depending on the application. The high temperature mechanical behavior of pure Mg had been investigated extensively [2–12]. Due to HCP structure, there were a few slip systems activated during deformation, a and c types dislocations, as well as (a + c) dislocation activated on both basal and non-basal planes [13–16]. Therefore, the poor ductility of Mg and Mg-alloys was attributed to the highly anisotropic dislocation slip behavior. Transmission electron microscope (TEM) observation of a 2%-elongated sample of AZ31 alloy revealed that the substantial dislocation slip is activated on non-basal planes by cross-slip from basal planes [17]. Twins can be found in AZ31 alloy during high temperature deformation, and the dynamic recovery (DRV) and dynamic recrystallization (DRX) play a significant role in reducing flow stress and raising ductility at elevated temperature [18]. However, only a few articles propose the mechanisms concerning
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twinning and recrystallization. The deformation mechanism and fracture features of the alloy during creep are still unclear. In the present study, by means of SEM and TEM observation, an investigation has been made into the microstructure evolution and the deformation features of AZ31 Mg-alloy during high temperature creep. 2. Experimental procedure Mg–Al–Zn alloy was used as the experimental materials, having a chemical composition of Mg–3Al–1Zn–0.2Mn in wt.%. The alloy was prepared in a mild steel crucible placed in a resistance furnace under an atmosphere of SF6 (1 vol.%) and CO2 (bal.) mixture, and mold cast into a block with 273 mm × 200 mm × 35 mm, then machined into the cylindrical specimens with a cross-section of 6 mm in diameter and 25 mm gauge length according to ASTM B557M. The uniaxial constant load tensile creep testing was conducted, using a PCM—L30M3 modal creep testing machine under applied stress of 50 MPa at 200 ◦ C, up to fracture. The strain value of creep was measured by means of linear variable differential transformer (LVDT) coupled with a data acquisition system. In order to understand the fracture features of the alloy during the high temperature creep and to provide a foundation, in further, for the design of a new heat resistance alloy, the experimental condition under an applied stress of 50 MPa at 200 ◦ C was chosen. To investigate the microstructure evolution regularity during creep, the alloy was cut into the sheet-like specimens with a cross-section of 3.4 mm × 1.5 mm and a gauge length of 25 mm. Before the creep testing was carried out, the wider surface of the specimens was polished, and the uniaxial constant load tensile tests under different condition were conducted in a high temperature microscope with the vacuum tensile apparatus (HM-4135 model) so that “in situ” observation of microstructure evolution in the alloy during creep may be made out. The strain data were measured with an extensometer. Tests were interrupted at different strain stages, and the specimens were cooled in the rate of 15 ◦ C min−1 to room temperature under load to preserve the dislocation structure. Microstructures of the alloy were observed by means of SEM and TEM. TEM foils sectioned are parallel to the applied stress axis, and thinned by conventional twins jet polishing technique using an electrolyte consisting of 5% HCIO4 , 35% butanol and 60% methanol at 253 K. A JEM-2000FXII transmission electron microscope at 200 kV was used for the morphology observation.
Fig. 1. Creep curve of AZ31 alloy under the condition of 50 MPa and 200 ◦ C.
initial strain (εo ) of creep. Then, the strain rate is reduced due to the increase of dislocation density as creep goes on. The period of steady state creep was about 70 h, the period being from 20 to 90 h as shown in Fig. 1, which displays the minimum strain rate measured as 1.45 × 10−6 (s−1 ). After crept for 100 h, the creep strain of the alloy was 0.03; the rupture time of the alloy under applied initial stress of 50 MPa at 200 ◦ C was 192 h. The microstructure of as-cast AZ31 alloy is shown in Fig. 2. The typical microstructure of the alloy consists of ␣-Mg matrix and fine particles that are homogeneously distributed in the matrix with volume fraction less than 5%. The fine particles are identified as Mg17 Al12 phase by XRD. Under the conditions of the applied stress of 50 MPa at 200 ◦ C, after the sheet-like specimen, laid in high temperature microscope with the vacuum tensile apparatus, was crept for 10 h, microstructure of the alloy is showed in Fig. 3. The inci¯ all of dislocations are dent beam direction is parallel to [1 2¯ 1 3], visible in g = 0 1¯ 1 1 diffraction condition. This indicates that significant amount of dislocation slips are activated in the matrix
3. Experimental results and analysis 3.1. Deformation features during initial creep The creep curve of AZ31 alloy was measured under the condition of an applied initial stress of 50 MPa at 200 ◦ C as shown in Fig. 1, representing that the alloy displays a large primary strain stage, including a transient strain and a shorter the reducing stage of creep rate. Transient strain is produced when load is applied at high temperature; at the same time dislocations multiply rapidly and fill in matrix of the alloy; this corresponds to the
Fig. 2. Structure of as-cast AZ31 alloy.
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Fig. 3. Dislocation configuration in the alloy crept for 10 h under applied stress ¯ of 50 MPa at 200 ◦ C; incident beam direction is parallel to [1 2¯ 1 3].
of alloy when the strain value reaches to 0.01. During tensile creep of the alloy, most of the activated dislocations are a dislocations slipping on the basal planes, having the Burgers vector of b = (1/3)2 1¯ 1¯ 0[19]. However, the distribution of the dislocations is different with the orientations of the grains supporting the tensile loading. Some of dislocations with a Burgers vector are activated on basal planes; another of dislocations having c or a + c Burgers vector are activated on non-basal planes, for example, activated on prismatic and pyramidal planes. After crept for 10 h, the dislocation configurations under the different diffraction conditions are shown in Fig. 4, the incident beam direction is parallel to [7 2¯ 5¯ 3]. This indicates that significant amount of the longer and straighter dislocations are activated in the matrix of the alloy. The characters of dislocations may be determined according to contrast analysis. All of dislocations are visible in g = 0 1 1¯ 1¯ (Fig. 4(a)) and g = 1¯ 1 0 3 (Fig. 4(b)) operation reflections, but, they are out of contrast in g = 1¯ 2 1¯ 2 as shown in Fig. 4(c). According to the g·b invisibility criterion, the dislocations of these ribbon-like contrasts have Burgers vector of (1/3)[1 2¯ 1 3]. Therefore, the longer and straighter dislocations are distinguished as a + c one, and may
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Fig. 5. Morphology of dislocations concentrated into cells after the alloy crept for 15 h under the condition of the applied stress of 50 MPa at 200 ◦ C. Incident beam direction is parallel to [0 1 1¯ 0].
slip on the (1¯ 2 1¯ 2) non-basal planes, which is defined as the pyramidal planes slip due to (1¯ 2 1¯ 2) plane laid in the pyramidal plane of HCP structure. The density of dislocations increases as creep goes on, which results in the reduction of strain rate. Under the conditions of the applied stress of 50 MPa at 200 ◦ C, after the sheet-like specimen crept for 15 h, the dislocations in the grain interiors are concentrated into the wall-like or cell-like arrangement as shown in Fig. 5. The incident beam direction is parallel to [0 1 1¯ 0], the wall-like or cell-like formed from dislocations re-arranged display contrast in g = 2¯ 1 1 0, which corresponds to the dynamic recovery of alloy during creep. 3.2. Deformation features during the steady state creep One of deformation features during the steady state creep is the twins occurred in the grain interiors. After the alloy crept for 25 h under applied stress of 50 MPa at 200 ◦ C, microstructures of the occurring twinning deformation are shown in Fig. 6. The twins patterns appear to be a rotations of the incident
Fig. 4. Contrast analysis of dislocation configuration of the alloy crept for 10 h under applied stress of 50 MPa at 200 ◦ C. Incident beam direction is parallel to ¯ (b) g = 1¯ 1 0 3; (c) g = 1¯ 2 1¯ 2. [7 2¯ 5¯ 3]: (a) g = 0 1 1¯ 1;
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Fig. 6. Twin morphology of the alloy crept for 25 h under applied stress of 50 MPa at 200 ◦ C: (a) twins in multiple parallel groups; (b) some of the twins frequently end in sharp points close to the boundaries as marked by white arrows, thin twin (H) cutting through the subgrains A, B, C and D, the serration twin boundaries as marked with black arrow.
electron beam, and somewhat more like diffraction loops in the poly-crystal regions [18] (the picture omitted); these occur entirely within individual crystal. The most evident feature of the microstructures is the thin twins often in the form of multiple parallel groups as shown in Fig. 6(a). Some of the twins frequently end in sharp points close to the boundaries as marked with white arrows in Fig. 6(b). The obvious evidence of DRV is the serrations appearance of the twin boundaries due to the absorbing some dislocation tangles as marked with black arrow in Fig. 6(b); this is the different from the matrix of alloy. In addition, there exist some dislocation tangles within the twins; these will be re-arranged into the walls or cells as creep goes on. A new phenomenon is the occurrence of very thin twins (H), which cut through the twins A, B, C, and end within the subgrain D as shown in Fig. 6(b). Significant amount of dislocations in alloy have been generated when the strain is 0.02, and then the ones are arranged into the cells and walls as shown in Fig. 5. As creep goes on, these cells encroach on each other to form the subgrains as shown in Fig. 7. Fig. 7(a) is a morphology of the substructure in which twin A possesses two parallel boundaries; dislocation cell B in twin A possesses a smooth sub-boundary which is considered to being protruded along the direction marked with arrow in Fig. 7(a); the cell can, in further, form the subgrain as creep continues as marked with C. The grains marked by A and C in Fig. 7(a) are the subgrains with different crystallogphy orientation. The subgrains in Fig. 7(a) and (b) are considered to be a
sub-structure produced during creep. These may act as the nucleation sites for promoting the occurrence of the recrystallization and grains growing up, so that, in further, is transformed into the microstructure in Fig. 8. Because a combination of the deforming dislocations and their concentration, and finally forming the subgrains is full accomplished during creep, the synthetical process is defined as the dynamic recrystallization. The grain sizes of the alloy prior to creep may be evaluated according to Fig. 2, about 80–100 m. It is found by means of the in situ observation under microscopy during high temperature creep that the grain size is changed as creep goes on, and gradually reduced until an equilibrium one is reached. After recrystallization, there is the different equilibrium grain size depending on the creep temperature. Under the applied stress of 50 MPa at different creep temperature, the different equilibrium grain sizes of the alloy are obtained as shown in Fig. 8. At same condition, there is a shorter creep fracture time for the sheet-like specimen than the cylindrical one due to the difference of the geometric shapes. The equilibrium grain size of the sheet-like specimen crept for 70 h at 200 ◦ C until fracture is 8–10 m as shown in Fig. 8(a), while the one after the alloy crept for 15 h increases to 30–40 m when temperature enhanced to 250 ◦ C at the applied same stress, the difference of the grain sizes is about several times as shown in Fig. 8(b) as compared with Fig. 8(a). Namely, the equilibrium size of the recrystallized grains increases with the enhancing creep temperature.
Fig. 7. Morphologies of subgrains originated from dynamic recrystallization in the alloy, crept for 50 h under applied stress of 50 MPa at 200 ◦ C: (a) co-existed morphology of twinning and subgrains, protruding of the subgrain boundary, in the cell B along the direction marked with arrow; (b) morphology of small subgrains.
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Fig. 8. SEM micrographs of the alloy crept up to fracture at different temperatures under applied stress of 50 MPa: (a) crept for 70 h at 200 ◦ C, displaying a feature along the boundaries fracture as marked by arrows; (b) crept for 15 h at 250 ◦ C.
Fig. 9. Schematic diagram of the relationship between the slip trace lines and applied stress direction.
3.3. Slip in grain interiors in the later period during the steady state creep It is found by means of the in situ observation under microscope that another deformation feature during creep is the slip activated in grain interiors. After the alloy is tensile crept for 10 h under the applied stress of 50 MPa at 250 ◦ C, the primary slip trace lines, in the first, appear in the grain interiors near the triple junction. Fig. 9 is a schematic diagram for showing the relationship between the slip trace lines and applied stress axis. The normal of the triple junction is perpendicular to the applied stress axis, the direction of the initial slip trace lines are also perpendicular to the applied stress axis. The areas of appearing the initial trace lines situate near the triple junction in which the normal of the triple junction is perpendicular to the applied stress
axis. As creep goes on, the areas of activating slip traces increase. Then, the slip traces are transformed into the wave morphology in some regions of the specimen as creep continues. After the alloy crept for 15 h until fracture, the number of slip trace lines increases and spreads in the grain interiors that seem to appear in the whole gauge length of the specimen. Morphology of the slip traces in the grain interiors is shown in Fig. 10; the direction of the slip traces with the wave form is about perpendicular to the applied stress axis as marked with arrows in Fig. 10(a), and the direction of secondary slip traces activated in some grain interiors appears to be at a θ angle relative to the initial slipping traces as marked with arrows in Fig. 10(b). Because the initial slip traces are first activated in the areas near the triple junction, it may be deducted that the stress concentration occurs easily in the areas near the triple junction, and the maximum value of the stress concentration occurs in the areas in which the normal of the triple junction is perpendicular to the applied stress axis. The slip in the grain interiors occurs when the maximum value of the stress concentration exceeds the resistance of the alloy, and releasing simultaneous the stress concentration. As creep goes on, the stress concentration occurs in the areas again, and the slip in the grain interiors is activated again, so that the number of the slip trace lines increase gradually as shown in Fig. 10(a) and (b). Whether the slip traces activated in the grain interiors, in the later period of the steady state creep, are corresponding to the dislocation slips on the basal and non-basal planes during the primary creep, this is to be discussed in further.
Fig. 10. Morphology of slip trace lines appearing on specimen surface in the alloy crept for 15 h under applied stress of 50 MPa at 250 ◦ C: (a) slip trace lines direction perpendicular to the applied stress axis; (b) multiple slip.
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Fig. 11. Morphology of the cracks formed and expanded along the boundaries in the tertiary stage of creep: (a) cracks formed in the triple junction; (b) cracks expanded along the boundary.
3.4. Continuous damage in the later period of creep An obvious feature of creep entering the tertiary stage is that the microcracks appear in the triple junction as shown in Fig. 11(a), and the normal of the triple junction is perpendicular to the applied stress axis yet. The areas appearing the microcracks are similar to the ones of the appearing initial slip traces. It is confirmed again from above that the stress concentration occurs easily in the areas near the triple junction. Because the stress concentration occurs easily in the location of the triple junction in which possesses the weakest bonding force in alloy, the microcracks occurs easily in the areas. As creep continues, the morphology of the microcrack expanded along the boundary is shown in Fig. 11(b), it is clearly shown that new wave slip traces appear in the grain interiors near the boundaries when the cracks expand as marked with arrow in Fig. 11(b). It may be deducted according to the morphology in Fig. 11(b) that the wave interfaces near the cavities is attributed to the activated slip lines originated from the cracks expanding along boundaries. The slip in the grain interiors occur again when the cracks expand along the boundaries, and the slip steps is kept on the viscous fracture interfaces as marked with arrows in Fig. 11(b). Namely, microstructure of the wave boundaries results from the slip steps formed from the microcracks expanding. The fracture feature of alloy crept up to failure is related to the binding energy between the grains, and also to the toughness or brittleness of the alloy. Due to a lower binding force existed in the grain boundary, therefore, the creep fracture of the alloy displays a feature of the cracks expanded viscous along the boundary as creep goes on as shown in Fig. 12. Some of the cavities appear round the grain when the cracks expand, and the viscous tearing strips are kept in the boundary regions round the grain as marked with arrows in Fig. 12. The bulky and continuous cavities in alloy result in such a structure that the grain is almost entirely separated from the other grains, and only parts of the tearing strips are interlinked with the other ones as shown in Fig. 12. This reduces the effective area of the alloy supporting the loading, results in the rapid enhancement of the strain rate. SEM observation indicates clearly that the homogeneous deformation occurs in the full gauge length in the sheet-like specimens, including both the slip in the grain interiors and the cracks viscous expanding along the boundary, which is consid-
Fig. 12. Morphology of the crack expanded viscous along the grain boundary.
ered to be the main reason for the rapid increase of the strain value in the later stage of creep. 4. Discussion 4.1. Compatibility stress and twinning deformation Some changes in microstructure occurred during creep has been proved by means of TEM and SEM observation. Dislocation slip is activated on the basal and non-basal planes in the initial stage of creep. In fact, AZ31 as a solid solution alloy in which dislocations slip is restricted by the presence of solute atmospheres, there will be an upper limiting stress at which the dislocations pull away from their atmospheres [20], the process is defined as solute drag creep. As the applied stress increase, in further, this may give rise to a transition from viscous slip to dislocations climb [21]. Then, dislocations are concentrated into the cells and walls, and in further the twinning and dynamic recrystallization occurs as creep goes on. This indicated that Mg-alloy with HCP structure displays a synthesizing deformation process during creep. For HCP structure materials, there are a few slip systems activated during creep [13–17,22–26], indicating that a strong
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heterogeneity of plastic deformation will occur in the alloy. The reports on the deformation mechanism of AZ31 alloy with HCP structure showed [17] that only a dislocations and a + c dislocations were activated on the basal and non-basal planes during the initial deformation. Because plastic yielding can occur in the Mg grain interiors at a very small stress of several MPa [17], namely, basal slip can occur without the influence of the compatibility stress, and large-grained alloys are expected to facilitate basal slip more easily than small-grained alloys. When basal slip plane in the grain interiors is approximately parallel to applied tensile stress axis, basal slip is stopped due to the resolved shear stress for basal slip nearly zero, but non-basal dislocation slip may be activated [17]. With the increase of strain to a certain level, twinning occurred as an additional deformation mechanism. The effects of compatibility stress on the grain boundaries have been investigated [27–30] in detail in various bi-crystal samples. The heterogeneity of plastic deformation is very strong in Mg-alloys, which may give rise to a large plastic compatibility stress near the grain boundaries. Whereas due to the heterogeneity deformation of HCP structure, a strong yield anisotropy in Mg prevents a large number of basal dislocations from slipping in the other orientation grains, which results in the rapid increase of the flow stress to sufficiently high level to activate twinning for the large-grained alloys as shown in Fig. 6. Thereby, it may be concluded that the twinning is an important deformation mechanism in polycrystalline alloy with HCP structure. Some experimental results [31] show that the deformation mechanism of the alloy at room temperature is the occurrence of dislocation slips and twins, thereinto, significant amount of dislocation slips are a dislocation activated on the basal and non-basal planes. During tensile creep in the range of applied temperatures and stresses, (a + c) dislocation may be activated on the non-basal planes, and there are still some twinning activated in the matrix of the alloy. But, compared with room temperature deformation, the probability of the occurred twinning deformation decreases gradually with the enhancement of creep temperature. Therefore, the activation of twinning has been known to be a necessary deformation mechanism for homogeneous deformation in the alloy. In the role of compatibility stress, the grains that cannot activate dislocation slip occur the twinning deformation in an appropriate orientation. The twin boundaries may both hinder dislocations motion and enhance the flow stress of the alloy. But the small change of the orientation in the twin interiors may promote the activation of dislocation slip again to correspond homogeneous deformation of the alloy. The supporting evidences of dislocation slip in the twin interiors are shown in Figs. 6 and 7 in [18]. Therefore, it may be deduced that the twin deformation may improve the ductility of some alloys with HCP structure. In the process of dynamic recovery during creep, the twin boundaries are changed into the serration due to absorbing movable dislocations, and the tangles in the grain interiors are concentrated into dislocations cells or walls. As creep progresses, the subgrains are gradually formed from the re-arrangement of the dislocation cells, walls and the protrusion of the serrated boundaries, which may serve as a core for promoting the occurrence of the dynamic recrystallization.
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4.2. Dynamic recrystallization It has been proven by SEM and TEM observation that the grain size changes during creep as shown in Figs. 7 and 8. That is attributed to the dynamic recrystallization during creep. In the temperature range of 200–250 ◦ C, the slip of the a and (a + c) dislocations activated on the basal and non-basal planes [32] is the controlling mechanism of the creep. The crossslip of an a dislocation by Frirdel–Escaig mechanism leads to a transition from a primary screw orientation to an edge orientation [33,34], and this edge dislocation can readily climb [35,36]. Dislocations originated from the cross-slipping and climbing are concentrated to form a low-angle boundary [17]. Continuous absorption of dislocations on the low-angle boundaries results in the formation of the serrated boundaries, the dislocation cells or walls are, in further, rearranged to form the high-angle subgrain boundaries as shown in Fig. 7, which provides the nucleation sites for DRX [18,25,26]. Then, the protrusion of subgrain boundaries as marked with arrow in Fig. 7(a), or the coalition of subgrain, indicates the occurrence of DRX. As creep goes on, the recrystallized grains grow to an equilibrium size, and there is a different equilibrium grain size after DRX during creep at different temperatures as shown in Fig. 8. Although the grain size after DRX is related to the applied stress and temperature [18,32,37] during creep, the difference of the grain size in Fig. 8(a) and (b) is attributed to the different creep temperature. The equilibrium grain size after DRX at 200 ◦ C is about 8–10 m as shown in Fig. 8(a), while that of DRX at 250 ◦ C is about 30–40 m as shown in Fig. 8(b). The finer grain structure from DRX may obviously improve the ductility of the alloy; therefore, it may be considered that the DRX is responsible for the rapid increase of the strain value in the later stage of creep. The slip traces appear within the recrystallized grains as shown in Fig. 9. If the occurrence of dislocations concentrating and twinning is considered to be the deformation mechanisms of the alloy in the steady state stage of creep, the slip in the grain interiors is the deformation mechanism in the later stage of steady state creep. Moreover, DRX during creep is responsible for improving the ductility of the alloy. 5. Conclusions 1. During the primary creep, the deformation features of the alloy are that significant amount of dislocations slip are activated on basal and non-basal planes; these ones are concentrated into the dislocation cells or walls, and further developed into the subgrains. Then, twinning occurs as an additional deformation mechanism during creep due to the plastic compatibility stress, which improves the ductility of the alloy. 2. The protrusion of subgrain boundaries indicates the occurrence of DRV. The finer grains in alloy result from DRX during steady state creep, while the equilibrium grain size after DRX is closely related to the creep temperature. The multiple slips in the grain interiors are considered to be the
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main deformation mechanism after DRX in the later stage of the steady state creep. 3. The initiation of the cracks on the grain boundaries indicates that creep enters the tertiary stage. As creep goes on, the cracks expanded viscous along the grain boundary result in the creep fracture of the alloy. 4. Dislocations slip and twinning are considered to be the main deformation mechanism of the alloy during creep. After DRX, the multiple slips in the grain interiors and the cracks expanded viscous along the grain boundary occur homogeneous in all length range of specimen, that, together with the twinning and dynamic recrystallization, results in the rapid increase of the strain rate in the later stage during creep. Acknowledgement Sponsorship of this research by the Natural Science Foundation of Liaoning Province in China under Grant no. 20022033 is gratefully acknowledged. References [1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [2] W.J.McG. Tegart, Acta Metall. 9 (1961) 614. [3] R.B. Jones, J.E. Harris, in: Proceeding of the Joint International Conference on Creep, the Institution of Mechanical Engineers, vol. 1, London, 1963, p. 1. [4] K. Milicka, J. Cadek, P. Rys, Acta Metall. 18 (1970) 1071. [5] S.S. Vagarali, T.G. Langdon, Acta Metall. 29 (1981) 1969. [6] L. Shi, D.O. Northwood, Acta Metall. Mater. 42 (1994) 871. [7] A. Orlova, K. Milicka, J. Cadek, Z. Metallkd. 63 (1972) 89. [8] S.S. Vagarali, T.G. Langdon, Acta Metall. 30 (1982) 1157. [9] W.K. Miller, in: H.G. Paris, W.H. Hunt (Eds.), Advances in Magnesium Allots and Composites, vol. 41, TMS, Warrendale, Pennsylvania, 1988. [10] E.M. Gutman, Y. Unigovski, M. Levkovich, Z. Koren, E. Anghion, M. Dangur, Mater. Sci. Eng. A 234–236 (1997) 880.
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Microstructure evolution and deformation features of AZ31 Mg-alloy during creep Sugui Tian a,∗ , Ling Wang a , Keun Yong Sohn b , Kyung Hyun Kim b , Yongbo Xu c , Zhuangqi Hu c a
School of Materials Science and Engineering, Shenyang University of Technology, No. 58 Xinghua South Street Tiexi District, Shenyang 110023, China b Korea Institute of Machinery and Materials, Changwon 641-010, South Korea c State Key Laboratory for Fatigue and Fracture of Materials, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China Received in revised form 3 October 2005; accepted 3 October 2005
Abstract By means of the measurement of the creep curve and the observation of SEM and transmission electron microscope (TEM), an investigation has been made into the microstructure evolution and deformation features of AZ31 Mg-alloy during high temperature creep. Results show that the deformation features of the alloy in the primary stage of creep are that significant amount of dislocation slips are activated on basal and non-basal planes, then these ones are concentrated into the dislocation cells or walls as creep goes on. At the same time, twinning occurs as an additional deformation mechanism in the role of the compatibility stress. During steady state creep, the dislocation cells are transformed into the subgrains, then, the protrusion and coalition of the sub-boundaries results in the occurrence of dynamic recovery (DRV). After the dynamic recrystallization (DRX), the multiple slips in the grain interiors are considered to be the main deformed mechanism in the later stage of the steady state creep. An obvious feature of creep entering the tertiary stage is that the cracks appear on the locations of the triple junction. As creep continues, the cracks are viscous expanded along the grain boundaries; this is taken for being the fracture mechanism of the alloy crept to failure. The multiple slips in the grain interiors and the cracks expanded viscous along the grain boundary occur in whole of specimens, that, together with the twins and dynamic recrystallization, is responsible for the rapid increase of the strain rate in the later stage during creep. © 2005 Elsevier B.V. All rights reserved. Keywords: AZ31 Mg-alloy; Creep; Dislocation slip; Twinning; Dynamic recrystallization; Deformation mechanism
1. Introduction Magnesium alloys are emerging as potentially good candidates for numerous applications, especially in the automotive industry. Their good properties, such as low density and high specific strength, make them promising replacements for other heavier materials like, for instance, steel, cast iron, and even aluminium [1]. However, the disadvantage of Mg-alloys is that they exhibit only limited ductility due to their HCP structure. Therefore, significant development effort is needed in order to widen the application of these materials. Magnesium alloys are being used to an increasing extent in applications where the components are subjected to elevated temperatures. Consequently, research is being focused on the
∗
Corresponding author. Tel.: +86 24 25694659; fax: +86 24 25691802. E-mail address: [email protected] (S. Tian).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.10.015
development of alloys able to withstand high stresses at temperature ranges of up to 180, 200 and 250 ◦ C depending on the application. The high temperature mechanical behavior of pure Mg had been investigated extensively [2–12]. Due to HCP structure, there were a few slip systems activated during deformation, a and c types dislocations, as well as (a + c) dislocation activated on both basal and non-basal planes [13–16]. Therefore, the poor ductility of Mg and Mg-alloys was attributed to the highly anisotropic dislocation slip behavior. Transmission electron microscope (TEM) observation of a 2%-elongated sample of AZ31 alloy revealed that the substantial dislocation slip is activated on non-basal planes by cross-slip from basal planes [17]. Twins can be found in AZ31 alloy during high temperature deformation, and the dynamic recovery (DRV) and dynamic recrystallization (DRX) play a significant role in reducing flow stress and raising ductility at elevated temperature [18]. However, only a few articles propose the mechanisms concerning
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twinning and recrystallization. The deformation mechanism and fracture features of the alloy during creep are still unclear. In the present study, by means of SEM and TEM observation, an investigation has been made into the microstructure evolution and the deformation features of AZ31 Mg-alloy during high temperature creep. 2. Experimental procedure Mg–Al–Zn alloy was used as the experimental materials, having a chemical composition of Mg–3Al–1Zn–0.2Mn in wt.%. The alloy was prepared in a mild steel crucible placed in a resistance furnace under an atmosphere of SF6 (1 vol.%) and CO2 (bal.) mixture, and mold cast into a block with 273 mm × 200 mm × 35 mm, then machined into the cylindrical specimens with a cross-section of 6 mm in diameter and 25 mm gauge length according to ASTM B557M. The uniaxial constant load tensile creep testing was conducted, using a PCM—L30M3 modal creep testing machine under applied stress of 50 MPa at 200 ◦ C, up to fracture. The strain value of creep was measured by means of linear variable differential transformer (LVDT) coupled with a data acquisition system. In order to understand the fracture features of the alloy during the high temperature creep and to provide a foundation, in further, for the design of a new heat resistance alloy, the experimental condition under an applied stress of 50 MPa at 200 ◦ C was chosen. To investigate the microstructure evolution regularity during creep, the alloy was cut into the sheet-like specimens with a cross-section of 3.4 mm × 1.5 mm and a gauge length of 25 mm. Before the creep testing was carried out, the wider surface of the specimens was polished, and the uniaxial constant load tensile tests under different condition were conducted in a high temperature microscope with the vacuum tensile apparatus (HM-4135 model) so that “in situ” observation of microstructure evolution in the alloy during creep may be made out. The strain data were measured with an extensometer. Tests were interrupted at different strain stages, and the specimens were cooled in the rate of 15 ◦ C min−1 to room temperature under load to preserve the dislocation structure. Microstructures of the alloy were observed by means of SEM and TEM. TEM foils sectioned are parallel to the applied stress axis, and thinned by conventional twins jet polishing technique using an electrolyte consisting of 5% HCIO4 , 35% butanol and 60% methanol at 253 K. A JEM-2000FXII transmission electron microscope at 200 kV was used for the morphology observation.
Fig. 1. Creep curve of AZ31 alloy under the condition of 50 MPa and 200 ◦ C.
initial strain (εo ) of creep. Then, the strain rate is reduced due to the increase of dislocation density as creep goes on. The period of steady state creep was about 70 h, the period being from 20 to 90 h as shown in Fig. 1, which displays the minimum strain rate measured as 1.45 × 10−6 (s−1 ). After crept for 100 h, the creep strain of the alloy was 0.03; the rupture time of the alloy under applied initial stress of 50 MPa at 200 ◦ C was 192 h. The microstructure of as-cast AZ31 alloy is shown in Fig. 2. The typical microstructure of the alloy consists of ␣-Mg matrix and fine particles that are homogeneously distributed in the matrix with volume fraction less than 5%. The fine particles are identified as Mg17 Al12 phase by XRD. Under the conditions of the applied stress of 50 MPa at 200 ◦ C, after the sheet-like specimen, laid in high temperature microscope with the vacuum tensile apparatus, was crept for 10 h, microstructure of the alloy is showed in Fig. 3. The inci¯ all of dislocations are dent beam direction is parallel to [1 2¯ 1 3], visible in g = 0 1¯ 1 1 diffraction condition. This indicates that significant amount of dislocation slips are activated in the matrix
3. Experimental results and analysis 3.1. Deformation features during initial creep The creep curve of AZ31 alloy was measured under the condition of an applied initial stress of 50 MPa at 200 ◦ C as shown in Fig. 1, representing that the alloy displays a large primary strain stage, including a transient strain and a shorter the reducing stage of creep rate. Transient strain is produced when load is applied at high temperature; at the same time dislocations multiply rapidly and fill in matrix of the alloy; this corresponds to the
Fig. 2. Structure of as-cast AZ31 alloy.
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Fig. 3. Dislocation configuration in the alloy crept for 10 h under applied stress ¯ of 50 MPa at 200 ◦ C; incident beam direction is parallel to [1 2¯ 1 3].
of alloy when the strain value reaches to 0.01. During tensile creep of the alloy, most of the activated dislocations are a dislocations slipping on the basal planes, having the Burgers vector of b = (1/3)2 1¯ 1¯ 0[19]. However, the distribution of the dislocations is different with the orientations of the grains supporting the tensile loading. Some of dislocations with a Burgers vector are activated on basal planes; another of dislocations having c or a + c Burgers vector are activated on non-basal planes, for example, activated on prismatic and pyramidal planes. After crept for 10 h, the dislocation configurations under the different diffraction conditions are shown in Fig. 4, the incident beam direction is parallel to [7 2¯ 5¯ 3]. This indicates that significant amount of the longer and straighter dislocations are activated in the matrix of the alloy. The characters of dislocations may be determined according to contrast analysis. All of dislocations are visible in g = 0 1 1¯ 1¯ (Fig. 4(a)) and g = 1¯ 1 0 3 (Fig. 4(b)) operation reflections, but, they are out of contrast in g = 1¯ 2 1¯ 2 as shown in Fig. 4(c). According to the g·b invisibility criterion, the dislocations of these ribbon-like contrasts have Burgers vector of (1/3)[1 2¯ 1 3]. Therefore, the longer and straighter dislocations are distinguished as a + c one, and may
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Fig. 5. Morphology of dislocations concentrated into cells after the alloy crept for 15 h under the condition of the applied stress of 50 MPa at 200 ◦ C. Incident beam direction is parallel to [0 1 1¯ 0].
slip on the (1¯ 2 1¯ 2) non-basal planes, which is defined as the pyramidal planes slip due to (1¯ 2 1¯ 2) plane laid in the pyramidal plane of HCP structure. The density of dislocations increases as creep goes on, which results in the reduction of strain rate. Under the conditions of the applied stress of 50 MPa at 200 ◦ C, after the sheet-like specimen crept for 15 h, the dislocations in the grain interiors are concentrated into the wall-like or cell-like arrangement as shown in Fig. 5. The incident beam direction is parallel to [0 1 1¯ 0], the wall-like or cell-like formed from dislocations re-arranged display contrast in g = 2¯ 1 1 0, which corresponds to the dynamic recovery of alloy during creep. 3.2. Deformation features during the steady state creep One of deformation features during the steady state creep is the twins occurred in the grain interiors. After the alloy crept for 25 h under applied stress of 50 MPa at 200 ◦ C, microstructures of the occurring twinning deformation are shown in Fig. 6. The twins patterns appear to be a rotations of the incident
Fig. 4. Contrast analysis of dislocation configuration of the alloy crept for 10 h under applied stress of 50 MPa at 200 ◦ C. Incident beam direction is parallel to ¯ (b) g = 1¯ 1 0 3; (c) g = 1¯ 2 1¯ 2. [7 2¯ 5¯ 3]: (a) g = 0 1 1¯ 1;
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Fig. 6. Twin morphology of the alloy crept for 25 h under applied stress of 50 MPa at 200 ◦ C: (a) twins in multiple parallel groups; (b) some of the twins frequently end in sharp points close to the boundaries as marked by white arrows, thin twin (H) cutting through the subgrains A, B, C and D, the serration twin boundaries as marked with black arrow.
electron beam, and somewhat more like diffraction loops in the poly-crystal regions [18] (the picture omitted); these occur entirely within individual crystal. The most evident feature of the microstructures is the thin twins often in the form of multiple parallel groups as shown in Fig. 6(a). Some of the twins frequently end in sharp points close to the boundaries as marked with white arrows in Fig. 6(b). The obvious evidence of DRV is the serrations appearance of the twin boundaries due to the absorbing some dislocation tangles as marked with black arrow in Fig. 6(b); this is the different from the matrix of alloy. In addition, there exist some dislocation tangles within the twins; these will be re-arranged into the walls or cells as creep goes on. A new phenomenon is the occurrence of very thin twins (H), which cut through the twins A, B, C, and end within the subgrain D as shown in Fig. 6(b). Significant amount of dislocations in alloy have been generated when the strain is 0.02, and then the ones are arranged into the cells and walls as shown in Fig. 5. As creep goes on, these cells encroach on each other to form the subgrains as shown in Fig. 7. Fig. 7(a) is a morphology of the substructure in which twin A possesses two parallel boundaries; dislocation cell B in twin A possesses a smooth sub-boundary which is considered to being protruded along the direction marked with arrow in Fig. 7(a); the cell can, in further, form the subgrain as creep continues as marked with C. The grains marked by A and C in Fig. 7(a) are the subgrains with different crystallogphy orientation. The subgrains in Fig. 7(a) and (b) are considered to be a
sub-structure produced during creep. These may act as the nucleation sites for promoting the occurrence of the recrystallization and grains growing up, so that, in further, is transformed into the microstructure in Fig. 8. Because a combination of the deforming dislocations and their concentration, and finally forming the subgrains is full accomplished during creep, the synthetical process is defined as the dynamic recrystallization. The grain sizes of the alloy prior to creep may be evaluated according to Fig. 2, about 80–100 m. It is found by means of the in situ observation under microscopy during high temperature creep that the grain size is changed as creep goes on, and gradually reduced until an equilibrium one is reached. After recrystallization, there is the different equilibrium grain size depending on the creep temperature. Under the applied stress of 50 MPa at different creep temperature, the different equilibrium grain sizes of the alloy are obtained as shown in Fig. 8. At same condition, there is a shorter creep fracture time for the sheet-like specimen than the cylindrical one due to the difference of the geometric shapes. The equilibrium grain size of the sheet-like specimen crept for 70 h at 200 ◦ C until fracture is 8–10 m as shown in Fig. 8(a), while the one after the alloy crept for 15 h increases to 30–40 m when temperature enhanced to 250 ◦ C at the applied same stress, the difference of the grain sizes is about several times as shown in Fig. 8(b) as compared with Fig. 8(a). Namely, the equilibrium size of the recrystallized grains increases with the enhancing creep temperature.
Fig. 7. Morphologies of subgrains originated from dynamic recrystallization in the alloy, crept for 50 h under applied stress of 50 MPa at 200 ◦ C: (a) co-existed morphology of twinning and subgrains, protruding of the subgrain boundary, in the cell B along the direction marked with arrow; (b) morphology of small subgrains.
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Fig. 8. SEM micrographs of the alloy crept up to fracture at different temperatures under applied stress of 50 MPa: (a) crept for 70 h at 200 ◦ C, displaying a feature along the boundaries fracture as marked by arrows; (b) crept for 15 h at 250 ◦ C.
Fig. 9. Schematic diagram of the relationship between the slip trace lines and applied stress direction.
3.3. Slip in grain interiors in the later period during the steady state creep It is found by means of the in situ observation under microscope that another deformation feature during creep is the slip activated in grain interiors. After the alloy is tensile crept for 10 h under the applied stress of 50 MPa at 250 ◦ C, the primary slip trace lines, in the first, appear in the grain interiors near the triple junction. Fig. 9 is a schematic diagram for showing the relationship between the slip trace lines and applied stress axis. The normal of the triple junction is perpendicular to the applied stress axis, the direction of the initial slip trace lines are also perpendicular to the applied stress axis. The areas of appearing the initial trace lines situate near the triple junction in which the normal of the triple junction is perpendicular to the applied stress
axis. As creep goes on, the areas of activating slip traces increase. Then, the slip traces are transformed into the wave morphology in some regions of the specimen as creep continues. After the alloy crept for 15 h until fracture, the number of slip trace lines increases and spreads in the grain interiors that seem to appear in the whole gauge length of the specimen. Morphology of the slip traces in the grain interiors is shown in Fig. 10; the direction of the slip traces with the wave form is about perpendicular to the applied stress axis as marked with arrows in Fig. 10(a), and the direction of secondary slip traces activated in some grain interiors appears to be at a θ angle relative to the initial slipping traces as marked with arrows in Fig. 10(b). Because the initial slip traces are first activated in the areas near the triple junction, it may be deducted that the stress concentration occurs easily in the areas near the triple junction, and the maximum value of the stress concentration occurs in the areas in which the normal of the triple junction is perpendicular to the applied stress axis. The slip in the grain interiors occurs when the maximum value of the stress concentration exceeds the resistance of the alloy, and releasing simultaneous the stress concentration. As creep goes on, the stress concentration occurs in the areas again, and the slip in the grain interiors is activated again, so that the number of the slip trace lines increase gradually as shown in Fig. 10(a) and (b). Whether the slip traces activated in the grain interiors, in the later period of the steady state creep, are corresponding to the dislocation slips on the basal and non-basal planes during the primary creep, this is to be discussed in further.
Fig. 10. Morphology of slip trace lines appearing on specimen surface in the alloy crept for 15 h under applied stress of 50 MPa at 250 ◦ C: (a) slip trace lines direction perpendicular to the applied stress axis; (b) multiple slip.
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Fig. 11. Morphology of the cracks formed and expanded along the boundaries in the tertiary stage of creep: (a) cracks formed in the triple junction; (b) cracks expanded along the boundary.
3.4. Continuous damage in the later period of creep An obvious feature of creep entering the tertiary stage is that the microcracks appear in the triple junction as shown in Fig. 11(a), and the normal of the triple junction is perpendicular to the applied stress axis yet. The areas appearing the microcracks are similar to the ones of the appearing initial slip traces. It is confirmed again from above that the stress concentration occurs easily in the areas near the triple junction. Because the stress concentration occurs easily in the location of the triple junction in which possesses the weakest bonding force in alloy, the microcracks occurs easily in the areas. As creep continues, the morphology of the microcrack expanded along the boundary is shown in Fig. 11(b), it is clearly shown that new wave slip traces appear in the grain interiors near the boundaries when the cracks expand as marked with arrow in Fig. 11(b). It may be deducted according to the morphology in Fig. 11(b) that the wave interfaces near the cavities is attributed to the activated slip lines originated from the cracks expanding along boundaries. The slip in the grain interiors occur again when the cracks expand along the boundaries, and the slip steps is kept on the viscous fracture interfaces as marked with arrows in Fig. 11(b). Namely, microstructure of the wave boundaries results from the slip steps formed from the microcracks expanding. The fracture feature of alloy crept up to failure is related to the binding energy between the grains, and also to the toughness or brittleness of the alloy. Due to a lower binding force existed in the grain boundary, therefore, the creep fracture of the alloy displays a feature of the cracks expanded viscous along the boundary as creep goes on as shown in Fig. 12. Some of the cavities appear round the grain when the cracks expand, and the viscous tearing strips are kept in the boundary regions round the grain as marked with arrows in Fig. 12. The bulky and continuous cavities in alloy result in such a structure that the grain is almost entirely separated from the other grains, and only parts of the tearing strips are interlinked with the other ones as shown in Fig. 12. This reduces the effective area of the alloy supporting the loading, results in the rapid enhancement of the strain rate. SEM observation indicates clearly that the homogeneous deformation occurs in the full gauge length in the sheet-like specimens, including both the slip in the grain interiors and the cracks viscous expanding along the boundary, which is consid-
Fig. 12. Morphology of the crack expanded viscous along the grain boundary.
ered to be the main reason for the rapid increase of the strain value in the later stage of creep. 4. Discussion 4.1. Compatibility stress and twinning deformation Some changes in microstructure occurred during creep has been proved by means of TEM and SEM observation. Dislocation slip is activated on the basal and non-basal planes in the initial stage of creep. In fact, AZ31 as a solid solution alloy in which dislocations slip is restricted by the presence of solute atmospheres, there will be an upper limiting stress at which the dislocations pull away from their atmospheres [20], the process is defined as solute drag creep. As the applied stress increase, in further, this may give rise to a transition from viscous slip to dislocations climb [21]. Then, dislocations are concentrated into the cells and walls, and in further the twinning and dynamic recrystallization occurs as creep goes on. This indicated that Mg-alloy with HCP structure displays a synthesizing deformation process during creep. For HCP structure materials, there are a few slip systems activated during creep [13–17,22–26], indicating that a strong
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heterogeneity of plastic deformation will occur in the alloy. The reports on the deformation mechanism of AZ31 alloy with HCP structure showed [17] that only a dislocations and a + c dislocations were activated on the basal and non-basal planes during the initial deformation. Because plastic yielding can occur in the Mg grain interiors at a very small stress of several MPa [17], namely, basal slip can occur without the influence of the compatibility stress, and large-grained alloys are expected to facilitate basal slip more easily than small-grained alloys. When basal slip plane in the grain interiors is approximately parallel to applied tensile stress axis, basal slip is stopped due to the resolved shear stress for basal slip nearly zero, but non-basal dislocation slip may be activated [17]. With the increase of strain to a certain level, twinning occurred as an additional deformation mechanism. The effects of compatibility stress on the grain boundaries have been investigated [27–30] in detail in various bi-crystal samples. The heterogeneity of plastic deformation is very strong in Mg-alloys, which may give rise to a large plastic compatibility stress near the grain boundaries. Whereas due to the heterogeneity deformation of HCP structure, a strong yield anisotropy in Mg prevents a large number of basal dislocations from slipping in the other orientation grains, which results in the rapid increase of the flow stress to sufficiently high level to activate twinning for the large-grained alloys as shown in Fig. 6. Thereby, it may be concluded that the twinning is an important deformation mechanism in polycrystalline alloy with HCP structure. Some experimental results [31] show that the deformation mechanism of the alloy at room temperature is the occurrence of dislocation slips and twins, thereinto, significant amount of dislocation slips are a dislocation activated on the basal and non-basal planes. During tensile creep in the range of applied temperatures and stresses, (a + c) dislocation may be activated on the non-basal planes, and there are still some twinning activated in the matrix of the alloy. But, compared with room temperature deformation, the probability of the occurred twinning deformation decreases gradually with the enhancement of creep temperature. Therefore, the activation of twinning has been known to be a necessary deformation mechanism for homogeneous deformation in the alloy. In the role of compatibility stress, the grains that cannot activate dislocation slip occur the twinning deformation in an appropriate orientation. The twin boundaries may both hinder dislocations motion and enhance the flow stress of the alloy. But the small change of the orientation in the twin interiors may promote the activation of dislocation slip again to correspond homogeneous deformation of the alloy. The supporting evidences of dislocation slip in the twin interiors are shown in Figs. 6 and 7 in [18]. Therefore, it may be deduced that the twin deformation may improve the ductility of some alloys with HCP structure. In the process of dynamic recovery during creep, the twin boundaries are changed into the serration due to absorbing movable dislocations, and the tangles in the grain interiors are concentrated into dislocations cells or walls. As creep progresses, the subgrains are gradually formed from the re-arrangement of the dislocation cells, walls and the protrusion of the serrated boundaries, which may serve as a core for promoting the occurrence of the dynamic recrystallization.
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4.2. Dynamic recrystallization It has been proven by SEM and TEM observation that the grain size changes during creep as shown in Figs. 7 and 8. That is attributed to the dynamic recrystallization during creep. In the temperature range of 200–250 ◦ C, the slip of the a and (a + c) dislocations activated on the basal and non-basal planes [32] is the controlling mechanism of the creep. The crossslip of an a dislocation by Frirdel–Escaig mechanism leads to a transition from a primary screw orientation to an edge orientation [33,34], and this edge dislocation can readily climb [35,36]. Dislocations originated from the cross-slipping and climbing are concentrated to form a low-angle boundary [17]. Continuous absorption of dislocations on the low-angle boundaries results in the formation of the serrated boundaries, the dislocation cells or walls are, in further, rearranged to form the high-angle subgrain boundaries as shown in Fig. 7, which provides the nucleation sites for DRX [18,25,26]. Then, the protrusion of subgrain boundaries as marked with arrow in Fig. 7(a), or the coalition of subgrain, indicates the occurrence of DRX. As creep goes on, the recrystallized grains grow to an equilibrium size, and there is a different equilibrium grain size after DRX during creep at different temperatures as shown in Fig. 8. Although the grain size after DRX is related to the applied stress and temperature [18,32,37] during creep, the difference of the grain size in Fig. 8(a) and (b) is attributed to the different creep temperature. The equilibrium grain size after DRX at 200 ◦ C is about 8–10 m as shown in Fig. 8(a), while that of DRX at 250 ◦ C is about 30–40 m as shown in Fig. 8(b). The finer grain structure from DRX may obviously improve the ductility of the alloy; therefore, it may be considered that the DRX is responsible for the rapid increase of the strain value in the later stage of creep. The slip traces appear within the recrystallized grains as shown in Fig. 9. If the occurrence of dislocations concentrating and twinning is considered to be the deformation mechanisms of the alloy in the steady state stage of creep, the slip in the grain interiors is the deformation mechanism in the later stage of steady state creep. Moreover, DRX during creep is responsible for improving the ductility of the alloy. 5. Conclusions 1. During the primary creep, the deformation features of the alloy are that significant amount of dislocations slip are activated on basal and non-basal planes; these ones are concentrated into the dislocation cells or walls, and further developed into the subgrains. Then, twinning occurs as an additional deformation mechanism during creep due to the plastic compatibility stress, which improves the ductility of the alloy. 2. The protrusion of subgrain boundaries indicates the occurrence of DRV. The finer grains in alloy result from DRX during steady state creep, while the equilibrium grain size after DRX is closely related to the creep temperature. The multiple slips in the grain interiors are considered to be the
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