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Superplastic behavior of an Al–Mg–Li alloy
Journal of Alloys and Compounds 487 (2009) 109–115
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Superplastic behavior of an Al–Mg–Li alloy Lingying Ye ∗ , Xinming Zhang ∗ , Dawei Zheng, Shengdan Liu, Jianguo Tang School of Materials Science and Engineering, Central South University, Changsha 410083, China
a r t i c l e
i n f o
Article history: Received 6 May 2009 Received in revised form 21 July 2009 Accepted 24 July 2009 Available online 3 August 2009 Keywords: Al–Mg–Li alloy Thermomechanical processing Grain refining Recrystallization Superplasticity
a b s t r a c t A novel thermomechanical processing has been developed for producing fine-grained Al–Mg–Li alloy 1420 sheet for superplasticity. The refined microstructure had layers of different grain structure through the normal direction of the sheet. The surface layer (SL) had fine, nearly equiaxed grains whereas the center layer (CL) had coarse, elongated grains. The superplastic properties and microstructural evolution of alloy 1420 with such microstructure were investigated using uniaxial tensile tests. It was shown that the refined alloy exhibited considerable elongations of 580–915% when deformed at 510–540 ◦ C and initial strain rates from 5 × 10−4 to 1 × 10−3 s−1 . Microstructure observation revealed that the alloy 1420 exhibited unusual cavitation behavior and extensive GBS occurred when tested at 510 ◦ C and above, however, the occurrence of GBS was not uniform. The effects of temperature and strain on grain size and shape, cavity size and volume fraction in the SL and CL were determined, respectively. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Currently, there is a great interest in developing ultra-fine, fully recrystallized grain structure in Al–Mg–Li alloys for superplasticity [1–4]. Moreover, these alloys with ultra-fine recrystallized grains show high service properties and enhanced workability. In contrast, these alloys show poor toughness and limited plasticity at room temperature due to extensive localization of plastic deformation in the unrecrystallized condition [5–7]. Thermomechanical processing (TMP) is usually used to produce fine-grained structure for superplastic aluminium alloys [8]. TMP uses overaging to develop large precipitates (∼1 m), and subsequent rolling to form deformation zones around these precipitates. Then fine-grained microstructures can be achieved by particle stimulated nucleation during the following recrystallization annealing. However, it was hard to use TMP consisting of cold or warm rolling to produce fine-grained Al–Mg–Li alloy sheets because of their poor workability at low temperatures that led to a premature fracture under working conditions [5,7]. Fortunately, a unique rolling method which used a high starting rolling temperature and a crossrolling method successfully prevented cracking of the sheets, and simultaneously fine grain structure of about 7 m was achieved by exploiting this rolling method in a novel TMP [9]. While studies of grain refinement and superplasticity of Al–Mg–Li alloy 1420 prepared by severe plastic deformation (SPD), such as equal-channel
∗ Corresponding authors. Tel.: +86 731 88830265; fax: +86 731 88830265. E-mail addresses: yelingying [email protected] (L. Ye), xmzhang [email protected] (X. Zhang). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.07.148
angular (ECA) pressing [1–4,10,11] and high-pressure torsion processing [12], were reported, there is no systematic investigation to be reported on the superplastic behavior of this alloy prepared by TMP. Therefore, the aim of the present work is to produce a microstructure that exhibits superplasticity by this novel TMP, attentions will be focused on the flow behavior and microstructural evolution of the thermomechanical processed Al–Mg–Li alloy 1420 sheets deformed at elevated temperature. 2. Experimental procedure The present study was carried out on the 8.2 mm thick hot rolled Al–Mg–Li alloy 1420 plates with chemical composition of Al–5.2% Mg–2.1% Li–0.12% Zr (wt.%). The plates were subjected to a novel TMP (Fig. 1(a)). The plates were first solution-treated at 475 ◦ C for 2 h, quenched in water and then overaging at 300 ◦ C/48 h + 400 ◦ C/4 h, subsequently deformed by a two-step rolling with 10–20% reduction per pass and a rolling speed of about 0.427 m/s as following: (i) The overaged plates were preheated at 400 ◦ C for 1 h, and then unidirectional rolled to reduction of 55% along the rolling direction (RD) of the hot rolled plate. (ii) Next, the plates were reheated at 340–380 ◦ C for 1 h, then rolled to the final thickness of about 1.7 mm with reduction of 54% by a cross-rolling. Here, crossrolling is referred to as a procedure in which the rolling direction is changed once by rotating the specimen by 90◦ about the normal direction, i.e., after preheating, the rolling direction is along the transverse direction (TD) of the original hot rolled plate. Finally the rolled sheets were recrystallized at 510 ◦ C for 0.5 h in a molten salt bath, then quenched in water. This novel TMP (Fig. 1(a)), rolling at a high starting temperature coupled with a cross-rolling after 55% rolled, produced the sheets with slightly edge fracture. At the same time, deformation zones around the precipitates formed by this rolling method were enough for the subsequent development of a fine grain structure by particle stimulated nucleation (PSN) due to the increased stored energy by lowering the re-heating temperature [9]. Following TMP (Fig. 1(a)), tensile samples were machined parallel to final rolling direction with gauge length
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Fig. 2. (a) Effect of temperature on true stress–true strain curves at an initial strain rate of 1 × 10−3 s−1 ; (b) effect of strain rate on true stress–true stain curves at 525 ◦ C.
Fig. 1. (a) Schematic illustration of TMP used to produce fine-grained Al–Mg–Li alloy 1420 sheet and (b) typical three-dimensional microstructure after TMP. The RD and TD were referred to the specimen frame of the final sheets.
of 1.2 and 7.7 m with a grain aspect ratio of 1.5 in the SL and CL, respectively. The proportion of CL, defined as the ratio of the thickness of CL to that of the whole sheet, was about 1/5. 3.2. Superplastic behavior
of 10 mm and gauge width of 6 mm. Tensile tests were performed at initial strain rates ranging from 1 × 10−2 to 5 × 10−4 s−1 in the temperature interval 450–570 ◦ C. A computer controlled MTS 810 testing machine equipped with a three zone split furnace was used. Temperature accuracy was within ± 1 ◦ C. Tensile samples were held for about 20 min before deformation. Specimens for optical metallography were annealed at 150 ◦ C for 13 h in order to decorate grain boundaries with second phase particles, then mechanically polished and etched by a special etching solution (1%HF + 4%HCl + 4%HNO3 + 15%H2 O in vol.%) and analyzed on a XJP-6A optical microscope. The average grain sizes in the rolling direction (dRD ), transverse direction (dTD ) and normal direction (dND ) were measured, respectively. Three-dimensional average grain size (dAVR ) was calculated according to dAVR = (dRD × dTD × dND )1/3 . Note that the RD and TD were referred to the specimen frame of the final sheets since they were produced by cross-rolling.
3. Results and discussion 3.1. Microstructure after TMP It is known that there usually exist obvious inhomogeneous distributions of recrystallized grain structure along the normal direction of the Al–Mg–Li alloy 1420 sheet after TMP [13,14]. The grain structure can be divided into three layers along the normal direction of the sheet. The two surface layers (SL) contain fine, relatively equiaxed grains whereas the center layer (CL) contains coarse, elongated grains. The three-dimensional grain structure produced by the novel TMP (Fig. 1(a)), shown in Fig. 1(b), exhibited an average grain diameter of 6.3 m with a grain aspect ratio
The typical true stress–true stain curves for the Al–Mg–Li alloy 1420 sheet deformed at an initial strain rate of 1 × 10−3 s−1 and in the temperature interval 450–570 ◦ C are shown in Fig. 2(a). At 450 ◦ C, the flow stress reaches a peak stress of about 24 MPa and then decreases sharply until fracture. With deformation temperature increasing to 480 ◦ C, the peak flow stress decreases drastically to about 7.5 MPa and a short steady state flow appears. At higher deformation temperatures of 510 ◦ C and above, the –ε curves suggest that a steady state of flow stress has been attained, whereby the stress varies a little with strain. The effect of strain rate on the nature of the true stress–true strain curves when deformed at 525 ◦ C and initial strain rates ranging from 5 × 10−4 to 1 × 10−2 s−1 is shown in Fig. 2(b). As expected, the flow stress increases with increasing strain rate. And, the rate of flow hardening is found to increase with an increase in strain rate. Furthermore, within the strain rate range investigated the stress–strain curves suggest the occurrence of steady state flow behavior at certain levels of strain. The span length of this strain, during which the steady state flow behavior appears, tends to increase with decreasing strain rate. The flow stress (taken at a strain of 0.4) as a function of strain rate is plotted on a double-logarithmic scale in Fig. 3(a). The stress–strain rate curves show a typical sigmoidal shape, especially at high deformation temperatures. The slope of the
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Fig. 3. Variation of (a) flow stress taken at ε = 0.4, (b) coefficient of strain rate sensitivity m taken at ε = 0.4 and (c) elongation to failure ı with strain rate.
ln(stress)–ln(strain rate) curve is known as the strain rate sensitivity index m. The values of m, obtained by applying regression analysis to the ln(stress)–ln(strain rate) data, are plotted as a function of initial strain rate in Fig. 3(b). When deformed at 510 ◦ C and above, the maximum m value of each temperature is found at an initial strain rate of 1 × 10−3 s−1 and with deformation temperature decreasing to 480 ◦ C and below it is found at an initial strain rate of 5 × 10−4 s−1 . It is also found that at a fixed initial strain rate, the m value tends to increase with increasing deformation temperature and the maximum m value of about 0.79 is found at 540 ◦ C and an initial strain rate of 1 × 10−3 s−1 . However, a maximum elongation to failure of about 915% is found at this initial strain rate but a lower deformation temperature of 525 ◦ C with a corresponding m valve of about 0.72. The variation of elongation to failure at different deformation temperatures with strain rate is shown in Fig. 3(c). It is seen that the refined alloy exhibits considerable elongations of 580–915% when deformed at 510–540 ◦ C and initial strain rates ranging from 5 × 10−4 to 1 × 10−3 s−1 . However, when expanding the strain rate to 5 × 10−3 s−1 and above, the tensile elongations decrease dramatically, especially at 1 × 10−2 s−1 , at which the alloy 1420 shows almost the same values in the temperature interval 450–540 ◦ C. The sample exhibiting the largest elongation is shown compared with that of an undeformed sample in Fig. 4 The variation in m value and elongation to failure at an initial strain rate of 1 × 10−3 s−1 as a function of deformation temperature is plotted in Fig. 5. The m value tends to increase monotonically with deformation temperature, whereas the maximum elongation to failure is achieved at 525 ◦ C and the elongation tends to decrease with further increasing temperature. The effect of strain on m value at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C is shown in Fig. 6,
it can be seen that the m value tends to decrease with increasing strain except that at a strain of about 1.5 it increases slightly under optimum superplastic deformation condition. 3.3. Microstructural evolution 3.3.1. Grain growth The microstructural evolution in both shoulder section and gauge section of the Al–Mg–Li alloy 1420 was investigated after superplastic deformation to failure and to different strain levels. The typical grain structures in the CLs of longitudinal sections of the tensile specimens deformed to failure at different temperatures and an initial strain rate of 1 × 10−3 s−1 is shown in Fig. 7. Average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the shoulder and gauge section of Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure at various temperatures are summarized in Table 1. The grain sizes and grain aspect ratios measured in the
Fig. 4. Superplastically deformed tensile specimen which fractured after 915% elongation, undeformed sample is shown for comparison.
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Fig. 5. Elongation to failure ı and coefficient of strain rate sensitivity m as functions of temperature at an initial strain rate of 1 × 10−3 s−1 .
SL and CL are listed, respectively. It is seen that in the grip section, i.e., under static annealing, grain structure both in SL and CL after TMP (in this work) was statically stable at 510 ◦ C and below. A pronounced grain growth occurred at 525 ◦ C and above. However, the values of GAR were almost not changed at all test temperatures. It is known that the stability of grain structure in the alloy 1420 is mainly due to the presence of high density of dispersed Al3 Zr and Al2 MgLi particles that effectively pin the grain boundaries. The
Fig. 6. Effect of strain on coefficient of strain rate sensitivity m at an initial strain rate of 1 × 10−3 s−1 .
Al3 Zr particles were reported not to be substantially dissolved even at 590 ◦ C and the Al2 MgLi particles are not completely dissolved at 490–500 ◦ C [14,15]. Thus the significant grain growth at 525 ◦ C and above might be due to the substantial dissolution of Al2 MgLi particles and partial dissolution of Al3 Zr particles. In the gauge section, superplastic deformation resulted in significant grain growth. And, when testing at 450–510 ◦ C, the grain sizes along the longitudinal direction increased faster than that along the normal direction, which led to elongation of grains. It can be evident from the increas-
Fig. 7. Microstructures of longitudinal section after superplastic deformation to failure at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 450 ◦ C, (b) 480 ◦ C, (c) 510 ◦ C, (d) 540 ◦ C.
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Fig. 8. Microstructural evolution of longitudinal section during superplastic deformation at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C: (a) ε = 0.5, (b) = 1.5, (c) ε = 1.9, (d) ε = 2.3.
ing values of GAR, as shown in Table 1, which indicates significant contribution of dislocation glide to total deformation [16]. However, when testing at 525 and 540 ◦ C, the resulting grain structures were nearly equiaxed and became more or less uniform through the normal direction, suggesting the elongated grains in the CL after TMP were gradually transformed to equiaxed grains during superplastic deformation. Fig. 8 shows the microstructures developed upon deformation with varying strains under the temperature and initial strain rate condition of 525 ◦ C and 1 × 10−3 s−1 . The micrographs shown in Fig. 8 were examined on the CL of the longitudinal section. The variations of average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in both the SL and CL of gauge section with strain are shown in Fig. 9. It is observed that the cavitations were initially formed at coarse, elongated grains in the CL at the early stage of deformation. With further straining, these cavitations grew and coalesced along tensile axis, and, new cavi-
tations started to form in the SL. These newly formed cavitations in the SL mainly located at triple points of relatively large grains and have irregular shapes. At the end stage of the deformation, the cavitations began to interlink along normal direction and led to the fracture of the sample. It is worth noting that superplastic deformation resulted in different grain growth behavior at different strain levels. It can be seen from Fig. 9 that at ε < 1.5, grain growth along the longitudinal direction was faster than that along normal direction, which resulted in increasing of GAR. However, at ε > 1.5, this trend was reversed, i.e., grain growth along the normal direction became faster than that along the longitudinal direction, resulted in equiaxed grain structures after deformation. The unusual variation of grain structure with strain could be a result of dynamic recrystallization that occurred at ε > 1.5. It is known that dynamic recrystallization during superplastic deformation could result in grain refinement which leads to the increasing of m value [16]. Though the dynamic recrystallization in the present study did not
Table 1 Average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the shoulder and gauge section of Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures. T (◦ C)
TMP 450 480 510 525 540 a
Total elongation (%)/the equivalent exposure time (s)
130/3400 345/6300 580/8150 915/12300 655/9550
Shoulder sectiona
Gauge sectiona
dL (m)
dN (m)
GAR
dL (m)
dN (m)
GAR
6.8/8.8 7.0/9.0 7.2/9.3 7.5/9.9 8.4/10.9 10.1/12.2
5.5/5.8 5.6/5.9 5.9/6.0 6.0/6.3 6.7/6.9 7.6/8.0
1.24/1.52 1.25/1.53 1.23/1.55 1.25/1.58 1.25/1.58 1.33/1.53
6.8/8.8 11.1/12.5 13.6/14.4 16.7/17.4 18.1/19.3 20.2/21.1
5.5/5.8 5.8/6.4 7.8/8.1 9.9/10.6 13.6/14.4 15.6/16.7
1.24/1.52 1.98/1.95 1.74/1.78 1.69/1.64 1.33/1.34 1.29/1.26
Numerator and denominator are grain sizes or grain aspect ratios measured in the surface layer and center layer, respectively.
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Fig. 9. Average grain size along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the gauge section of Al–Mg–Li alloy 1420 as a function of strain when deformed at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C.
result in grain refinement, but a significant reduction in grain aspect ratio was observed. So the slight increase of m value (Fig. 6) at a strain around 1.5 could be associated with the dynamic recrystallization. 3.3.2. Cavitation The average cavity size A and cavity volume fraction V in the Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures are summarized in Table 2. The characteristics of cavitations were examined in the area 5 mm apart from the fracture surface. The majority of large cavities were observed to grow along the tensile direction and to exhibit an irregular shape suggesting plasticity-controlled cavity growth [16]. Two typical micrographs of cavities resulting from superplastic deformation at 480 and 525 ◦ C are shown in Fig. 10. It is noted that the average cavity size tends to decrease when deformed at 510 ◦ C whereas porosity volume fraction increases monotonously with increasing testing temperature. It is important to note that the reduction in average cavity size is mainly caused by the increasing number of cavities. Raising the deformation temperature from 480 to 510 ◦ C led to the superplastic elongation increasing from 345% to 580%, this is mainly due to the increasing contributions of grain boundary sliding (GBS) to the total elongation at a higher temperature. Besides contribution to superplastic elongation, extensive GBS at 510 ◦ C also resulted in increasing number of cavities. Though the porosity volume fraction also increased with increasing deformation temperature, its increasing speed was not as fast as that of numbers of cavities until at 525 ◦ C and above. 3.3.3. Grain boundary sliding It is interesting to note that some grains in the failure samples that tested at 510 ◦ C and above were revealed without etching after mechanically polishing, as shown in Fig. 11. This is suggested to be a result of GBS and the role of GBS tends to increase with increasing test temperature. The examined areas shown in Fig. 11 cover the
Fig. 10. Cavitations of longitudinal section after superplastic deformation to failure at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 480 ◦ C, (b) 525 ◦ C.
whole thickness of the longitudinal section of the failure samples. It is also found that the GBS does not occur uniformly during the superplastic deformation. At all testing temperatures, the occurrence of GBS is more severe in the center location than that in the edge location of sample. Furthermore, at 510 ◦ C, it can be seen from Fig. 11(a) that among the grains there are some areas in which the grains are not revealed, which suggests there GBS occurs in a cooperative manner through the shift of grain groups as a unit [17]. It has been reported [17,18] that cooperative grain boundary sliding (CGBS), i.e., sliding of grain groups, is an important mechanism of superplastic flow in fine-grained materials. Zelin and Mukherjee reported that the average size of these grain groups in a Pb–62%Sn eutectic alloy was about four to eight units of the average matrix grain size [19]. In the present study, it is about four to six times of the matrix grain size (Fig. 11(a)). Kaibyshev et al. [17] investigated the CGBS behavior in a fine-grained Al–Mg alloy found that at low temperatures, CGBS occurred along favorably oriented grain boundary surfaces, such as the grain boundaries with the maximum values of Schmid factor, and, increasing strain or testing temperature could led to transition from the sliding of grain groups to the sliding of individual grains. However, in the present study, it is noted that the
Table 2 Average cavity size A, and cavity volume fraction V, for the Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures.
Total elongation (%)/the equivalent exposure time (s) A (m) V (%)
450 ◦ C
480 ◦ C
510 ◦ C
525 ◦ C
540 ◦ C
130/3400 3.4 2.9
345/6300 5.8 5.8
580/8150 5.3 14.2
915/12300 6.5 14.8
655/9550 7.0 15.6
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val 450–540 ◦ C with the maximum total elongation of about 915% when deformed at a temperature of 525 ◦ C and an initial strain rate of 1 × 10−3 s−1 with a corresponding m valve of about 0.72. During superplastic deformation, the coarse, elongated grains in the CL gradually transformed to equiaxed grains and the cavitations were found to nucleate and grow in the CL firstly. With further straining, these cavitations grew and coalesced along tensile axis, and, new cavitations started to form in the SL. At the end stage of the deformation, the cavitations began to interlink along normal direction and led to the fracture of the sample. Microstructure observation revealed that extensive GBS occurred when tested at 510 ◦ C and above, however, the occurrence of GBS was not uniform. Acknowledgement This work was supported by the International Science and Technology Cooperation Program of China (no. 2008DFR50260) and the Major State Basic Research Development Program of China (no. 2005CB623706). References
Fig. 11. Micrographs showing non-uniform grain boundary sliding occurs during superplastic deformation at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 510 ◦ C, (b) 540 ◦ C.
uniformity of GBS was increased when test temperature increasing to 540 ◦ C (Fig. 11(b)) but increasing strain at 510 ◦ C did not result in a significant influence on the homogeneity of GBS. 4. Conclusions In summary, the present study demonstrates that the Al–Mg–Li alloy 1420 sheet prepared by the novel TMP contained layers of different microstructures along the normal direction of the sheet. The two SL contained fine, relatively equiaxed grains whereas the CL contained coarse, elongated grains. The alloy 1420 with such microstructure exhibited superplasticity in the temperature inter-
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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Superplastic behavior of an Al–Mg–Li alloy Lingying Ye ∗ , Xinming Zhang ∗ , Dawei Zheng, Shengdan Liu, Jianguo Tang School of Materials Science and Engineering, Central South University, Changsha 410083, China
a r t i c l e
i n f o
Article history: Received 6 May 2009 Received in revised form 21 July 2009 Accepted 24 July 2009 Available online 3 August 2009 Keywords: Al–Mg–Li alloy Thermomechanical processing Grain refining Recrystallization Superplasticity
a b s t r a c t A novel thermomechanical processing has been developed for producing fine-grained Al–Mg–Li alloy 1420 sheet for superplasticity. The refined microstructure had layers of different grain structure through the normal direction of the sheet. The surface layer (SL) had fine, nearly equiaxed grains whereas the center layer (CL) had coarse, elongated grains. The superplastic properties and microstructural evolution of alloy 1420 with such microstructure were investigated using uniaxial tensile tests. It was shown that the refined alloy exhibited considerable elongations of 580–915% when deformed at 510–540 ◦ C and initial strain rates from 5 × 10−4 to 1 × 10−3 s−1 . Microstructure observation revealed that the alloy 1420 exhibited unusual cavitation behavior and extensive GBS occurred when tested at 510 ◦ C and above, however, the occurrence of GBS was not uniform. The effects of temperature and strain on grain size and shape, cavity size and volume fraction in the SL and CL were determined, respectively. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Currently, there is a great interest in developing ultra-fine, fully recrystallized grain structure in Al–Mg–Li alloys for superplasticity [1–4]. Moreover, these alloys with ultra-fine recrystallized grains show high service properties and enhanced workability. In contrast, these alloys show poor toughness and limited plasticity at room temperature due to extensive localization of plastic deformation in the unrecrystallized condition [5–7]. Thermomechanical processing (TMP) is usually used to produce fine-grained structure for superplastic aluminium alloys [8]. TMP uses overaging to develop large precipitates (∼1 m), and subsequent rolling to form deformation zones around these precipitates. Then fine-grained microstructures can be achieved by particle stimulated nucleation during the following recrystallization annealing. However, it was hard to use TMP consisting of cold or warm rolling to produce fine-grained Al–Mg–Li alloy sheets because of their poor workability at low temperatures that led to a premature fracture under working conditions [5,7]. Fortunately, a unique rolling method which used a high starting rolling temperature and a crossrolling method successfully prevented cracking of the sheets, and simultaneously fine grain structure of about 7 m was achieved by exploiting this rolling method in a novel TMP [9]. While studies of grain refinement and superplasticity of Al–Mg–Li alloy 1420 prepared by severe plastic deformation (SPD), such as equal-channel
∗ Corresponding authors. Tel.: +86 731 88830265; fax: +86 731 88830265. E-mail addresses: yelingying [email protected] (L. Ye), xmzhang [email protected] (X. Zhang). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.07.148
angular (ECA) pressing [1–4,10,11] and high-pressure torsion processing [12], were reported, there is no systematic investigation to be reported on the superplastic behavior of this alloy prepared by TMP. Therefore, the aim of the present work is to produce a microstructure that exhibits superplasticity by this novel TMP, attentions will be focused on the flow behavior and microstructural evolution of the thermomechanical processed Al–Mg–Li alloy 1420 sheets deformed at elevated temperature. 2. Experimental procedure The present study was carried out on the 8.2 mm thick hot rolled Al–Mg–Li alloy 1420 plates with chemical composition of Al–5.2% Mg–2.1% Li–0.12% Zr (wt.%). The plates were subjected to a novel TMP (Fig. 1(a)). The plates were first solution-treated at 475 ◦ C for 2 h, quenched in water and then overaging at 300 ◦ C/48 h + 400 ◦ C/4 h, subsequently deformed by a two-step rolling with 10–20% reduction per pass and a rolling speed of about 0.427 m/s as following: (i) The overaged plates were preheated at 400 ◦ C for 1 h, and then unidirectional rolled to reduction of 55% along the rolling direction (RD) of the hot rolled plate. (ii) Next, the plates were reheated at 340–380 ◦ C for 1 h, then rolled to the final thickness of about 1.7 mm with reduction of 54% by a cross-rolling. Here, crossrolling is referred to as a procedure in which the rolling direction is changed once by rotating the specimen by 90◦ about the normal direction, i.e., after preheating, the rolling direction is along the transverse direction (TD) of the original hot rolled plate. Finally the rolled sheets were recrystallized at 510 ◦ C for 0.5 h in a molten salt bath, then quenched in water. This novel TMP (Fig. 1(a)), rolling at a high starting temperature coupled with a cross-rolling after 55% rolled, produced the sheets with slightly edge fracture. At the same time, deformation zones around the precipitates formed by this rolling method were enough for the subsequent development of a fine grain structure by particle stimulated nucleation (PSN) due to the increased stored energy by lowering the re-heating temperature [9]. Following TMP (Fig. 1(a)), tensile samples were machined parallel to final rolling direction with gauge length
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Fig. 2. (a) Effect of temperature on true stress–true strain curves at an initial strain rate of 1 × 10−3 s−1 ; (b) effect of strain rate on true stress–true stain curves at 525 ◦ C.
Fig. 1. (a) Schematic illustration of TMP used to produce fine-grained Al–Mg–Li alloy 1420 sheet and (b) typical three-dimensional microstructure after TMP. The RD and TD were referred to the specimen frame of the final sheets.
of 1.2 and 7.7 m with a grain aspect ratio of 1.5 in the SL and CL, respectively. The proportion of CL, defined as the ratio of the thickness of CL to that of the whole sheet, was about 1/5. 3.2. Superplastic behavior
of 10 mm and gauge width of 6 mm. Tensile tests were performed at initial strain rates ranging from 1 × 10−2 to 5 × 10−4 s−1 in the temperature interval 450–570 ◦ C. A computer controlled MTS 810 testing machine equipped with a three zone split furnace was used. Temperature accuracy was within ± 1 ◦ C. Tensile samples were held for about 20 min before deformation. Specimens for optical metallography were annealed at 150 ◦ C for 13 h in order to decorate grain boundaries with second phase particles, then mechanically polished and etched by a special etching solution (1%HF + 4%HCl + 4%HNO3 + 15%H2 O in vol.%) and analyzed on a XJP-6A optical microscope. The average grain sizes in the rolling direction (dRD ), transverse direction (dTD ) and normal direction (dND ) were measured, respectively. Three-dimensional average grain size (dAVR ) was calculated according to dAVR = (dRD × dTD × dND )1/3 . Note that the RD and TD were referred to the specimen frame of the final sheets since they were produced by cross-rolling.
3. Results and discussion 3.1. Microstructure after TMP It is known that there usually exist obvious inhomogeneous distributions of recrystallized grain structure along the normal direction of the Al–Mg–Li alloy 1420 sheet after TMP [13,14]. The grain structure can be divided into three layers along the normal direction of the sheet. The two surface layers (SL) contain fine, relatively equiaxed grains whereas the center layer (CL) contains coarse, elongated grains. The three-dimensional grain structure produced by the novel TMP (Fig. 1(a)), shown in Fig. 1(b), exhibited an average grain diameter of 6.3 m with a grain aspect ratio
The typical true stress–true stain curves for the Al–Mg–Li alloy 1420 sheet deformed at an initial strain rate of 1 × 10−3 s−1 and in the temperature interval 450–570 ◦ C are shown in Fig. 2(a). At 450 ◦ C, the flow stress reaches a peak stress of about 24 MPa and then decreases sharply until fracture. With deformation temperature increasing to 480 ◦ C, the peak flow stress decreases drastically to about 7.5 MPa and a short steady state flow appears. At higher deformation temperatures of 510 ◦ C and above, the –ε curves suggest that a steady state of flow stress has been attained, whereby the stress varies a little with strain. The effect of strain rate on the nature of the true stress–true strain curves when deformed at 525 ◦ C and initial strain rates ranging from 5 × 10−4 to 1 × 10−2 s−1 is shown in Fig. 2(b). As expected, the flow stress increases with increasing strain rate. And, the rate of flow hardening is found to increase with an increase in strain rate. Furthermore, within the strain rate range investigated the stress–strain curves suggest the occurrence of steady state flow behavior at certain levels of strain. The span length of this strain, during which the steady state flow behavior appears, tends to increase with decreasing strain rate. The flow stress (taken at a strain of 0.4) as a function of strain rate is plotted on a double-logarithmic scale in Fig. 3(a). The stress–strain rate curves show a typical sigmoidal shape, especially at high deformation temperatures. The slope of the
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Fig. 3. Variation of (a) flow stress taken at ε = 0.4, (b) coefficient of strain rate sensitivity m taken at ε = 0.4 and (c) elongation to failure ı with strain rate.
ln(stress)–ln(strain rate) curve is known as the strain rate sensitivity index m. The values of m, obtained by applying regression analysis to the ln(stress)–ln(strain rate) data, are plotted as a function of initial strain rate in Fig. 3(b). When deformed at 510 ◦ C and above, the maximum m value of each temperature is found at an initial strain rate of 1 × 10−3 s−1 and with deformation temperature decreasing to 480 ◦ C and below it is found at an initial strain rate of 5 × 10−4 s−1 . It is also found that at a fixed initial strain rate, the m value tends to increase with increasing deformation temperature and the maximum m value of about 0.79 is found at 540 ◦ C and an initial strain rate of 1 × 10−3 s−1 . However, a maximum elongation to failure of about 915% is found at this initial strain rate but a lower deformation temperature of 525 ◦ C with a corresponding m valve of about 0.72. The variation of elongation to failure at different deformation temperatures with strain rate is shown in Fig. 3(c). It is seen that the refined alloy exhibits considerable elongations of 580–915% when deformed at 510–540 ◦ C and initial strain rates ranging from 5 × 10−4 to 1 × 10−3 s−1 . However, when expanding the strain rate to 5 × 10−3 s−1 and above, the tensile elongations decrease dramatically, especially at 1 × 10−2 s−1 , at which the alloy 1420 shows almost the same values in the temperature interval 450–540 ◦ C. The sample exhibiting the largest elongation is shown compared with that of an undeformed sample in Fig. 4 The variation in m value and elongation to failure at an initial strain rate of 1 × 10−3 s−1 as a function of deformation temperature is plotted in Fig. 5. The m value tends to increase monotonically with deformation temperature, whereas the maximum elongation to failure is achieved at 525 ◦ C and the elongation tends to decrease with further increasing temperature. The effect of strain on m value at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C is shown in Fig. 6,
it can be seen that the m value tends to decrease with increasing strain except that at a strain of about 1.5 it increases slightly under optimum superplastic deformation condition. 3.3. Microstructural evolution 3.3.1. Grain growth The microstructural evolution in both shoulder section and gauge section of the Al–Mg–Li alloy 1420 was investigated after superplastic deformation to failure and to different strain levels. The typical grain structures in the CLs of longitudinal sections of the tensile specimens deformed to failure at different temperatures and an initial strain rate of 1 × 10−3 s−1 is shown in Fig. 7. Average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the shoulder and gauge section of Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure at various temperatures are summarized in Table 1. The grain sizes and grain aspect ratios measured in the
Fig. 4. Superplastically deformed tensile specimen which fractured after 915% elongation, undeformed sample is shown for comparison.
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Fig. 5. Elongation to failure ı and coefficient of strain rate sensitivity m as functions of temperature at an initial strain rate of 1 × 10−3 s−1 .
SL and CL are listed, respectively. It is seen that in the grip section, i.e., under static annealing, grain structure both in SL and CL after TMP (in this work) was statically stable at 510 ◦ C and below. A pronounced grain growth occurred at 525 ◦ C and above. However, the values of GAR were almost not changed at all test temperatures. It is known that the stability of grain structure in the alloy 1420 is mainly due to the presence of high density of dispersed Al3 Zr and Al2 MgLi particles that effectively pin the grain boundaries. The
Fig. 6. Effect of strain on coefficient of strain rate sensitivity m at an initial strain rate of 1 × 10−3 s−1 .
Al3 Zr particles were reported not to be substantially dissolved even at 590 ◦ C and the Al2 MgLi particles are not completely dissolved at 490–500 ◦ C [14,15]. Thus the significant grain growth at 525 ◦ C and above might be due to the substantial dissolution of Al2 MgLi particles and partial dissolution of Al3 Zr particles. In the gauge section, superplastic deformation resulted in significant grain growth. And, when testing at 450–510 ◦ C, the grain sizes along the longitudinal direction increased faster than that along the normal direction, which led to elongation of grains. It can be evident from the increas-
Fig. 7. Microstructures of longitudinal section after superplastic deformation to failure at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 450 ◦ C, (b) 480 ◦ C, (c) 510 ◦ C, (d) 540 ◦ C.
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Fig. 8. Microstructural evolution of longitudinal section during superplastic deformation at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C: (a) ε = 0.5, (b) = 1.5, (c) ε = 1.9, (d) ε = 2.3.
ing values of GAR, as shown in Table 1, which indicates significant contribution of dislocation glide to total deformation [16]. However, when testing at 525 and 540 ◦ C, the resulting grain structures were nearly equiaxed and became more or less uniform through the normal direction, suggesting the elongated grains in the CL after TMP were gradually transformed to equiaxed grains during superplastic deformation. Fig. 8 shows the microstructures developed upon deformation with varying strains under the temperature and initial strain rate condition of 525 ◦ C and 1 × 10−3 s−1 . The micrographs shown in Fig. 8 were examined on the CL of the longitudinal section. The variations of average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in both the SL and CL of gauge section with strain are shown in Fig. 9. It is observed that the cavitations were initially formed at coarse, elongated grains in the CL at the early stage of deformation. With further straining, these cavitations grew and coalesced along tensile axis, and, new cavi-
tations started to form in the SL. These newly formed cavitations in the SL mainly located at triple points of relatively large grains and have irregular shapes. At the end stage of the deformation, the cavitations began to interlink along normal direction and led to the fracture of the sample. It is worth noting that superplastic deformation resulted in different grain growth behavior at different strain levels. It can be seen from Fig. 9 that at ε < 1.5, grain growth along the longitudinal direction was faster than that along normal direction, which resulted in increasing of GAR. However, at ε > 1.5, this trend was reversed, i.e., grain growth along the normal direction became faster than that along the longitudinal direction, resulted in equiaxed grain structures after deformation. The unusual variation of grain structure with strain could be a result of dynamic recrystallization that occurred at ε > 1.5. It is known that dynamic recrystallization during superplastic deformation could result in grain refinement which leads to the increasing of m value [16]. Though the dynamic recrystallization in the present study did not
Table 1 Average grain sizes along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the shoulder and gauge section of Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures. T (◦ C)
TMP 450 480 510 525 540 a
Total elongation (%)/the equivalent exposure time (s)
130/3400 345/6300 580/8150 915/12300 655/9550
Shoulder sectiona
Gauge sectiona
dL (m)
dN (m)
GAR
dL (m)
dN (m)
GAR
6.8/8.8 7.0/9.0 7.2/9.3 7.5/9.9 8.4/10.9 10.1/12.2
5.5/5.8 5.6/5.9 5.9/6.0 6.0/6.3 6.7/6.9 7.6/8.0
1.24/1.52 1.25/1.53 1.23/1.55 1.25/1.58 1.25/1.58 1.33/1.53
6.8/8.8 11.1/12.5 13.6/14.4 16.7/17.4 18.1/19.3 20.2/21.1
5.5/5.8 5.8/6.4 7.8/8.1 9.9/10.6 13.6/14.4 15.6/16.7
1.24/1.52 1.98/1.95 1.74/1.78 1.69/1.64 1.33/1.34 1.29/1.26
Numerator and denominator are grain sizes or grain aspect ratios measured in the surface layer and center layer, respectively.
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Fig. 9. Average grain size along longitudinal direction dL , normal direction dN and the grain aspect ratio GAR in the gauge section of Al–Mg–Li alloy 1420 as a function of strain when deformed at an initial strain rate of 1 × 10−3 s−1 and 525 ◦ C.
result in grain refinement, but a significant reduction in grain aspect ratio was observed. So the slight increase of m value (Fig. 6) at a strain around 1.5 could be associated with the dynamic recrystallization. 3.3.2. Cavitation The average cavity size A and cavity volume fraction V in the Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures are summarized in Table 2. The characteristics of cavitations were examined in the area 5 mm apart from the fracture surface. The majority of large cavities were observed to grow along the tensile direction and to exhibit an irregular shape suggesting plasticity-controlled cavity growth [16]. Two typical micrographs of cavities resulting from superplastic deformation at 480 and 525 ◦ C are shown in Fig. 10. It is noted that the average cavity size tends to decrease when deformed at 510 ◦ C whereas porosity volume fraction increases monotonously with increasing testing temperature. It is important to note that the reduction in average cavity size is mainly caused by the increasing number of cavities. Raising the deformation temperature from 480 to 510 ◦ C led to the superplastic elongation increasing from 345% to 580%, this is mainly due to the increasing contributions of grain boundary sliding (GBS) to the total elongation at a higher temperature. Besides contribution to superplastic elongation, extensive GBS at 510 ◦ C also resulted in increasing number of cavities. Though the porosity volume fraction also increased with increasing deformation temperature, its increasing speed was not as fast as that of numbers of cavities until at 525 ◦ C and above. 3.3.3. Grain boundary sliding It is interesting to note that some grains in the failure samples that tested at 510 ◦ C and above were revealed without etching after mechanically polishing, as shown in Fig. 11. This is suggested to be a result of GBS and the role of GBS tends to increase with increasing test temperature. The examined areas shown in Fig. 11 cover the
Fig. 10. Cavitations of longitudinal section after superplastic deformation to failure at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 480 ◦ C, (b) 525 ◦ C.
whole thickness of the longitudinal section of the failure samples. It is also found that the GBS does not occur uniformly during the superplastic deformation. At all testing temperatures, the occurrence of GBS is more severe in the center location than that in the edge location of sample. Furthermore, at 510 ◦ C, it can be seen from Fig. 11(a) that among the grains there are some areas in which the grains are not revealed, which suggests there GBS occurs in a cooperative manner through the shift of grain groups as a unit [17]. It has been reported [17,18] that cooperative grain boundary sliding (CGBS), i.e., sliding of grain groups, is an important mechanism of superplastic flow in fine-grained materials. Zelin and Mukherjee reported that the average size of these grain groups in a Pb–62%Sn eutectic alloy was about four to eight units of the average matrix grain size [19]. In the present study, it is about four to six times of the matrix grain size (Fig. 11(a)). Kaibyshev et al. [17] investigated the CGBS behavior in a fine-grained Al–Mg alloy found that at low temperatures, CGBS occurred along favorably oriented grain boundary surfaces, such as the grain boundaries with the maximum values of Schmid factor, and, increasing strain or testing temperature could led to transition from the sliding of grain groups to the sliding of individual grains. However, in the present study, it is noted that the
Table 2 Average cavity size A, and cavity volume fraction V, for the Al–Mg–Li alloy 1420 strained at an initial strain rate of 1 × 10−3 s−1 up to failure under various temperatures.
Total elongation (%)/the equivalent exposure time (s) A (m) V (%)
450 ◦ C
480 ◦ C
510 ◦ C
525 ◦ C
540 ◦ C
130/3400 3.4 2.9
345/6300 5.8 5.8
580/8150 5.3 14.2
915/12300 6.5 14.8
655/9550 7.0 15.6
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val 450–540 ◦ C with the maximum total elongation of about 915% when deformed at a temperature of 525 ◦ C and an initial strain rate of 1 × 10−3 s−1 with a corresponding m valve of about 0.72. During superplastic deformation, the coarse, elongated grains in the CL gradually transformed to equiaxed grains and the cavitations were found to nucleate and grow in the CL firstly. With further straining, these cavitations grew and coalesced along tensile axis, and, new cavitations started to form in the SL. At the end stage of the deformation, the cavitations began to interlink along normal direction and led to the fracture of the sample. Microstructure observation revealed that extensive GBS occurred when tested at 510 ◦ C and above, however, the occurrence of GBS was not uniform. Acknowledgement This work was supported by the International Science and Technology Cooperation Program of China (no. 2008DFR50260) and the Major State Basic Research Development Program of China (no. 2005CB623706). References
Fig. 11. Micrographs showing non-uniform grain boundary sliding occurs during superplastic deformation at an initial strain rate of 1 × 10−3 s−1 and different temperatures: (a) 510 ◦ C, (b) 540 ◦ C.
uniformity of GBS was increased when test temperature increasing to 540 ◦ C (Fig. 11(b)) but increasing strain at 510 ◦ C did not result in a significant influence on the homogeneity of GBS. 4. Conclusions In summary, the present study demonstrates that the Al–Mg–Li alloy 1420 sheet prepared by the novel TMP contained layers of different microstructures along the normal direction of the sheet. The two SL contained fine, relatively equiaxed grains whereas the CL contained coarse, elongated grains. The alloy 1420 with such microstructure exhibited superplasticity in the temperature inter-
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