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The role of microstructures on the strengthening mechanisms of a thermomechanically processed 2091 Al–Li alloy
Materials Science and Engineering A284 (2000) 14 – 24 www.elsevier.com/locate/msea
The role of microstructures on the strengthening mechanisms of a thermomechanically processed 2091 Al–Li alloy Heon-Joo Kim a,*, Mitsuo Niinomi b a
Department of Metallurgical Engineering, Pukyong National Uni6ersity, San 100 Yongdang-dong, Nam-ku, Pusan 608 -739, South Korea b Department of Production System Engineering, Toyohashi Uni6ersity of Technology, Tempaku, Toyohashi 441, Japan Received 27 October 1999; received in revised form 9 February 2000
Abstract The relationship between tensile properties and microstructural parameters are discussed. Grain size, degree of recrystallization, size and distribution of precipitates were varied by thermomechanical processing. Tensile tests were performed and the microstructures were characterized using transmission electron microscopy (TEM). Quantitative analysis was carried out to explain the yield strength in connection with the change in the microstructural parameters. The primary strengthening precipitates of the under aged alloy were d% (Al3Li) and S% (Al2CuMg). The greater the warm rolling reduction ratio and the lower warm rolling temperature give a higher strength. Closely distributed fine S% precipitates were effective in improving the yield strength. The increment in yield strength by S% precipitation can be explained quantitatively by the Orowan’s mechanism. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Tensile properties of 2091 Al–Li alloy; Thermomechanical processing (TMP); d% (Al3Li) and S% (Al2CuMg) precipitates; Quantitative analysis of microstructural parameters; Orowan’s mechanism
1. Introduction Much attention has been paid to develop aluminum– lithium alloys that have the combination of lower density with increased stiffness and strength. The reduction in density and increase in stiffness are attributed to the lithium addition to aluminum which is a desirable property for aircraft structural applications. Consequently, numerous studies have been carried out on various mechanical and microstructural characteristics of these alloys [1 – 5]. A thermomechanical processing (TMP) [6] is a useful step in the grain refinement of Al – Li alloys [7]. Fine grains which are achieved by TMP, are effective in reducing the local stress concentration of coarse planar slips at grain boundaries by shortening the dislocation pile-up length in Al – Li alloys [8]. Depending on the variables in TMP, the degree of recrystallization as well as grain size can be controlled; * Corresponding author. Tel.: +82-51-6201480; fax: + 82-516250718. E-mail address: [email protected] (H.-J. Kim)
microstructures of completely recrystallized, partially recrystallized or completely unrecrystallized grains with different grain sizes can be obtained by the proper selection of temperature and deformation mode in TMP. The deficiencies in the Al–Li alloys can thus be improved through the use of microstructural control using TMP techniques. The mechanical properties of Al–Li alloys vary with the degree of recrystallization. According to the papers by Starke et al. [7] and Jata et al. [9], the unrecrystallized microstructure of 2020 Al– Li alloy increases the strain to failure compared with the recrystallized microstructure. The objective of this study is to analyze the effect of microstructural features on the strength of 2091 Al–Li alloy. A quantitative study was carried out, especially focusing on explaining the change in yield strength due to d% (Al3Li) and S% (Al2CuMg) precipitates strengthening mechanism. Changes of the microstructure in this alloy, i.e. grain size, degree of recrystallization and size, spacing, vol.% of precipitates were accomplished by changes in the TMP.
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2. Experimental procedures The material used in this study is 2091 Al –Li alloy, whose chemical composition in wt.% is 2.15Li, 2.10Cu, 1.49Mg, 0.12Zr and balance of Al. The iron, silicon and sodium are 0.06 wt.%, 0.03 wt.% and 3 ppm, respectively. The original ingots (1300×350 × 40 mm) of alloy were homogenized, then 5 mm planed off from both surfaces, processed with different TMPs and rolled to a final thickness of 6 mm plates. The alloy was processed with various TMPs as shown in Fig. 1 schematically. Detailed processes of the TMP are as follows. 1. The warm rolling reduction ratio was varied at 30, 45, 60 and 75% for a constant warm rolling temperature of 623 K. 2. The warm rolling temperature was varied to 573, 623 and 723 K for a constant warm rolling reduction ratio of 45%. 3. The cold rolling reduction ratio prior to solution treatment was varied at 0, 20 and 30%, after warm rolling at the condition of 623 K, 45%. 4. Cold rolling reduction ratio was varied at 0, 20
15
and 40%, after a constant 45% warm rolling condition at 623 K and annealing treatments, 673 K, 2 h. The maximum reduction ratio of cold rolling was 30% in case (c), compared with 40% in case of (d). This was due to the occurrence of edge cracking when an annealing step was not used before cold rolling. The TMP processes were designed to control the microstructural parameter of the precipitates, grain size, subgrain size and degree of recrystallization. Also 3% cold rolling just before artificial aging was conducted to promote an even distribution of the fine S% precipitates within all of the specimens. The underaging treatment, in which toughness of the alloy is increased by lesser precipitation at the grain boundaries [10], was finally carried out in the present TMP. The purpose of the design of these TMP processes can be found [11]. Tensile specimens were machined to plate type with the dimension of 2 mm thickness, 4 mm gage width and 20 mm gage length, with the axis parallel to the rolling direction of the plates. Tensile tests were performed at room temperature at a crosshead speed of 8.3× 10 − 6 m s − 1 using an Instron type machine. The 0.2% yield strength, ultimate tensile strength, Young’s modulus and fracture strain were measured from the tensile test results. Microstructures of the alloys were characterized using an optical microscopy and transmission electron microscopy, TEM. Polished samples for optical microscope observation were etched by immersion in Keller’s reagent. Thin foil for TEM were prepared with a twin-jet Tenupol apparatus using a 20% methyl acetate solution at 253 K and a potential difference of 20 V. Thickness of the TEM thin foils was measured according to the contamination spot method [12] for minimizing the error in measuring the microstructural parameters concerning precipitates. A computer image analyzer was also used to measure the shape and distribution of microstructural parameters.
3. Results
3.1. Microstructural obser6ation
Fig. 1. Schematic drawing of thermomechanical processings applied in this study.
Table 1 shows the observation results of representative samples by an optical microscope. The microstructures of all the samples treated with different warm rolling reduction ratios were completely unrecrystallized after aging. The grains became more fine as the warm rolling temperature decreased. But elongated grains were observed at higher temperatures.
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Table 1 Variation of degree of recrystallization and grain size with processing condition Designation
(a) Warm rolling reduction ratio change (At constant temperature, 623 K) (b) Warm rolling temperature change (At constant ratio, 45%) (c) Cold rolling reduction ratio change (At constant warm rolling condition, 623 K, 45%) (d) Annealing and cold rolling (At constant annealing condition, 673 K, 2 h)
Degree of recrystallization (%)
30% 75% 573 K 723 K 20% 30% 20% 40%
0 0 0 0 0 7.5 0 5.1
The grains were slightly coarser when the cold rolling reduction ratio was 20% in comparison with the 30% ratio. Owing to the effect of severe cold rolling (30%), some grains were recrystallized at the grain boundaries after successive processing. When annealing was carried out before cold rolling, some precipitates formed at the grain boundaries. Change in grain size was appreciable with cold rolling and no annealing. The grains were more refined in the 20% ratio sample than in the case of 40%, as compared to those with cold rolling without annealing. Some recrystallized grains also nucleated at the grain boundaries when the ratio was 40%. d% (Al3Li) and S% (Al2CuMg) precipitates which are primary strengthening precipitates, were confirmed by TEM analysis in the under aged 2091 Al – Li alloy. While this composite d% precipitates were present in some TMP conditions, needle or cylindrical S% precipitates were found under all conditions.
Grain size Longtudinal (mm)
Transverse (mm)
Short transverse (mm)
51.6 45.2 38.1 65.1 50.0 50.0 35.1 41.3
42.9 22.6 30.2 31.7 28.6 26.2 14.9 22.2
54.0 16.7 20.2 23.8 23.8 19.0 15.1 21.5
3.2. Effect of warm rolling reduction ratio on tensile property Fig. 2 exhibits the TEM micrographs of subgrains when the warm rolling reduction ratios are 30 and 75% for warm rolling temperature of 623 K, in which spherical d% and needle S% are precipitated. The microstructural parameters measured on TEM micrographs for each condition are listed in Table 2. Tensile properties for different warm rolling reduction ratios are shown in Table 3. While the change of 0.2% yield strength is negligible, tensile strength and fracture strain are increased as the warm rolling reduction ratio increases. A prominent increase in tensile strength and fracture strain is shown when the ratio is changed to 60%. At this condition, drastic changes followed, i.e. decrease in the diameter and average spacing between the nearest neighbor in d% precipitates and increase in the length and average spacing between the nearest neighbor in S% precipitates.
Fig. 2. Transmission electron microscopy (TEM) showing subgrain structure, d% and S% precipitates when warm rolling reduction ratio changed.
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Table 2 Effect of warm rolling reduction ratio on subgrain size and microstructural parameters Warm rolling reduction ratio (%)
Subgrain size (mm)
d% phase Volume fraction (%)
30 45 60 75
2.5 2.2 1.9 1.8
0.70 0.73 0.70 0.67
Diameter (nm) 32 31 27 26
Average spacing (nm) 77.7 79.0 64.4 59.1
S% phase
30 45 60 75
2.5 2.2 1.9 1.8
Volume fraction (%)
Length (nm)
0.23 0.22 0.20 0.20
70.5 83.2 98.5 111
Width (nm)
Average spacing (nm)
6 5 5 5
67.7 54.6 91.4 97.4
Table 3 Effect of warm rolling reduction ratio on tensile properties Warm rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
30 45 60 75
388 393 394 396
402 410 451 490
74.5 75.4 76.7 78.6
0.097 0.102 0.129 0.132
a
Yield strength is the stress at 0.2% offset strain.
Table 4 Effect of warm rolling temperature on subgrain size and microstructural parameters Warm rolling temperature (K)
Subgrain size (mm)
d% phase Volume fraction (%)
723 623 573
3.1 2.2 2.1
0.75 0.73 1.44
Diameter (nm) 33 31 15
Average spacing (nm) 67.9 79.0 30.5
S% phase
723 623 573
3.1 2.2 2.1
Volume fraction (%)
Length (nm)
0.22 0.22 0.15
79.9 83.2 125
Width (nm) 6 5 5
Average spacing (nm) 55.8 54.6 38.5
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Table 5 Effect of warm rolling temperature on tensile properties Warm rolling temperature (K)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
723 623 573
359 388 394
401 409 429
75.9 75.6 76.1
0.037 0.052 0.062
a
Yield strength is the stress at 0.2% offset strain.
Fig. 3. Transmission electron microscopy (TEM) micrographs showing subgrain structure, d% and S% precipitates when cold rolling reduction ratio changed.
Table 6 Effect of cold rolling reduction ratio before solution treatment on subgrain size and microstructural parameters Cold rolling reduction ratio (%)
0 20 30
Subgrain size (mm)
2.2 1.7 1.9
d% phase Volume fraction (%)
Diameter (nm)
Average spacing (nm)
0.73 1.38 1.71
31 19 22
79.0 65.9 55.7
Volume fraction (%)
Length (nm)
Width (nm)
0.22 0.11 0.16
83.2 60.1 57.8
S% phase
0 20 30
2.2 1.7 1.9
5 5 5
Average spacing (nm) 54.6 79.6 63.4
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smallest subgrain size and the most refined d% precipitates.
3.3. Effect of warm rolling temperature on tensile property Microstructural parameters for warm rolling temperature of 573, 623 and 723 K are shown in Table 4 when the warm rolling reduction ratio is kept constant at 45%. Table 5 shows the variation of the tensile properties as a function of the warm rolling temperature. Most of the tensile properties increase as the temperature de-creases. At a constant warm rolling reduction ratio, strength is maximum at the condition of the
3.4. Effect of cold rolling reduction ratio on tensile property Fig. 3 and Table 6 show TEM micrographs and microstructural parameters as the cold rolling reduction ratio changes, respectively. The effect of cold rolling ratio before solution heat treatment on tensile properties is shown in Table 7. All
Table 7 Effect of cold rolling reduction ratio before solution treatment on tensile properties Cold rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
0 20 30
389 394 402
419 447 488
76.6 75.6 75.0
0.049 0.062 0.067
a
Yield strength is the stress at 0.2% offset strain.
Table 8 Effect of cold rolling reduction ratio after annealing on subgrain size and microstructural parameters Cold rolling reduction ratio (%)
0 20 40
Subgrain size (mm)
2.3 1.8 2.1
d% phase Volume fraction (%)
Diameter (nm)
Average spacing (nm)
0.98 0.96 0.70
24 21 27
38.3 43.8 51.9
Volume fraction (%)
Length (nm)
Width (nm)
0.12 0.11 0.20
59.6 75.0 98.5
S% phase
0 20 40
2.3 1.8 2.1
5 5 5
Average spacing (nm) 63.8 101.2 55.3
Table 9 Effect of cold rolling reduction ratio after annealing on tensile properties Cold rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
0 20 40
363 387 417
398 449 485
75.7 76.4 78.1
0.028 0.039 0.046
a
Yield strength is the stress at 0.2% offset strain.
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The resolved shear yield stress, ty, of the present alloy can be described as; ty = tAl + DtG.B. + DtS.S. + DtP
Fig. 4. Relationship between subgrain diameter, d − 1/2, and 0.2% yeild strength, s0.2, of pure aluminum.
(1)
where tAl is the resolved shear yield stress of the Al matrix, and tG.B., tS.S. and tP are the increments of the resolved shear stress associated with grain refinement, solid solution and precipitation strengthenings, respectively. In the following discussion, the contribution of each strengthening mechanism to the 0.2% yield strength of the present alloy is determined based on theoretical analysis or equations. Finally, each constituent strength is added up to compare with the experimental overall yield strength of the 2091 alloy.
4.1. Grain refinement strengthening Yield strength of the alloy related to grain size can be described by the following Hall–Petch equation [13];
Fig. 5. Schematic drawing of calculating solid solution strengthening.
the tensile properties except elastic modulus increase as the cold rolling reduction ratio increases.
3.5. Successi6e effect of annealing and cold rolling on tensile property Microstructural parameters for annealing and different cold rolling reduction ratios are shown in Table 8. Some precipitates appeared at grain boundaries though precipitates free zones were not detected. Table 9 shows the variation of the mechanical properties as a function of cold rolling reduction ratio, after annealing treatments. Tensile strength shows the maximum value at the most severe cold reduction ratio, 40%, that is, at the condition of the maximum volume fraction and the longest length of S% precipitates. All the values are, however, smaller than those without annealing treatments. These low values are due to the appearance of precipitates at grain boundaries.
4. Discussion The strengthening mechanism in the present alloy will be theoretically discussed in the following sections.
1 s0.2 = s0 + k (2)
d where s0 is the friction stress, k is the constant and d is the subgrain diameter. Fig. 4 shows the 0.2% yield strength of pure aluminum against d − 1/2. A linear relation is observed between s0.2 and d − 1/2 for pure aluminum. The data of pure Al reported by Wong et al. [14] are used in calculating the yield strength increment of the Al matrix in the following section. The slope, i.e. the increment of yield strength associated with grain refinement is 21 MPa mm − 1/2.
4.2. Solid solution strengthening The degree of solid solution strengthening for precipitation hardening aluminum alloys is, generally, evaluated when an alloy is in the as-quenched condition. The degree of solid solution strengthening was, therefore, modified by assuming that it decreased logarithmically with aging time as schematically shown in Fig. 5. The effect of solid solution strengthening could be assumed to weaken gradually and completely disappear by reaching at 24 h aging time because this alloy has exhibited peak hardness value at this aging time. The aging time adopted in this study was 10 h. The degree of solid solution strengthening for as-solutionized alloy was calculated by subtracting s0.2 and an increment in strength caused by grain refinement for pure aluminum which is the result of Section 4.1 from s0.2 of as-quenched alloy. The degree of non-modified solid solution strengthening of the alloy is from 140 to 150 MPa, while the modified one, i.e. the degree of solid solution strengthening is from 35 to 50 MPa. It was postulated that the rate of precipitation was constant
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during aging, and also it was the same in all the materials used.
(c) Strengthening due to misfit strain between matrix and d%;
4.3. Precipitation strengthening
Dto =
The main precipitates in the present alloy were found to be d% and S% from TEM observation. Strengthening mechanisms of Al – Li binary alloys by d% has been well defined, and good agreement between experimental and calculated values of strength has been reported [15]. The strengthening due to S% has not been examined in detail. The contributions of d% and S% to s0.2 of the present alloy will be discussed focussing on the strengthening mechanisms due to S%.
4.3.1. Strengthening by d% precipitates Dislocations move by shearing d%, because d% is coherent with the matrix. Noble et al. [15] have summarized that the strengthening of Al – Li type alloy by d% is associated with the following four mechanisms; (a) Strengthening due to the creation of antiphase boundary by dislocation pairs; g 4f −f (3) Dt0 = 2b p where Dt0 is the increment in resolved shear strength due to the creation of an antiphase boundary by dislocation pairs. (b) Strengthening due to the difference in shear modulus between the matrix and d%; DG 3DG r 3 · 0.8 − 0.143 ln · r · f DtG = 2 · 4p Gb b (4) where DtG is the increment in resolved shear strength due to the difference in shear modulus between matrix and d%, and DG is the difference in shear modulus between matrix and d%.
'
'
'!
"
(3G o 3 r f)
b
(5)
where Dto is the increment in resolved shear strength caused by misfit strain, G is the shear modulus of matrix (30 GPa), r is the effective radius of d%, o is the misfit strain (0.7 × 10 − 3) due to the precipitation of d%, f is the volume fraction of the d% precipitates and b is the Burgers’ vector (0.2864 nm). (d) Strengthening due to shearing of d% by dislocations; DtS =
6 f · gS · p r
(6)
where DtS is the increment in resolved shear strength due to shearing of d% by dislocations and gS is the surface energy of the matrix (0.2 J m − 2). The equation of s= Mt where M is the Taylor factor was used for converting resolved shear stress, t to tensile stress, s. Although M is a function of the processing parameter due to texture [16–18], M= 3.07 which had been proposed by Decker [13] and generally used for f.c.c. metals. The calculated values of the increment in tensile stress, according to the above-mentioned four strengthening mechanisms, are shown in Fig. 6. The increment in strength due to the creation of antiphase boundaries by dislocation pairs is greatest, approaching 75–107 MPa while those due to the misfit strain between d% precipitates and matrix, and shearing of d% precipitates by dislocations are very small.
4.3.2. Strengthening by S% precipitates The strengthening mechanism of commercial Al–Li alloys by S% precipitates has remained as unclear. The chemical strengthening due to the shearing of S% precipitates by dislocations and the strengthening due to the dislocation movement according to the Orowan’s by-pass mechanism were applied for explaining the strengthening of the present alloy by S% precipitates. 4.3.2.1. Chemical strengthening. The increments in resolved shear stress caused by weak and strong precipitates, Dtch,w, and Dtch,s are given by the following equations, respectively[19]; For weak precipitate, Dtch,w =
Fig. 6. Strengthening increments associated with d% precipitation obtained from various strengthening mechanisms.
'
4g 3bf GZD
(7)
where gi is the surface energy of precipitates, G (= 0.5Gb 2) is the line tension of dislocation, f is the
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stress by screw dislocation is given by the following Orowan’s equation modified by Ashby [22];
Z r0 DtSC = 2pD(1− y) Gb ln
Fig. 7. Comparison between experimental S% phase precipitation strengthening and theoretical one calculated by Orowan’s mechanism on yield strength.
Fig. 8. Relationship between precipitation strengthening increment for yield strength and S% phase parameters.
(9)
where D is the effective distance between S% precipitates, y is the Poisson’s ratio (0.33) and r0 (= 2b) is the radius of the dislocation core. The values of 1.41 times of the measured width and distance of S% precipitates were substituted for Z and D, respectively, because S% precipitated on (210)100 which crossed the matrix slip plane, (111), by 45° [23]. The measured increment in strength caused by S% precipitates can be given by subtracting the theoretically calculated strength increments caused by grain refinement, solid solution and d% precipitates from measured s0.2 of the aged alloy. Both modified and nonmodified strength increments by solid solution strengthening were taken into account in evaluating the measured increment in strength by S% precipitates. The calculated strength increment by S% precipitates which is based on the chemical strengthening mechanism of weak and strong precipitates and the measured one are compared. The calculated values are much smaller than the measured ones. The strength increment by S% precipitates, therefore, can not be explained by the chemical strengthening mechanism. The calculated strength increments by S% precipitates according to the Orowan’s mechanism are compared with measured ones. The calculated strength increment with linear correction of solid solution strengthening and measured one are shown in Fig. 7, in which Taylor factor M=3.07 was used. The calculated values are in good agreement with the measured values.
volume fraction of the S% precipitates, Z is the effective width of S% precipitates and D is the length of S% precipitates. For strong precipitate, Dtch,s = 1.62gi
'
f ZD
(8)
The value of gi was measured to be 1.18 J m − 2 using the reversion technique reported by Hung et al. [20] in the present study.
4.3.2.2. Strengthening due to Orowan’s mechanism. According to the report by Miura et al. [21], screw dislocations are predominant in the plastic deformation of Al–Li type alloys. The increment in resolved shear
Fig. 9. Comparison between measured and theoretical yield strength.
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Therefore, the increment in strength caused by S% precipitation can be explained by Orowan’s mechanism based on the above theoretical analysis. The increment of yield strength associated with S% precipitation by Orowan’s mechanism is assumed to be from 100–200 MPa.
4.4. Analysis strengths
of
measured
and
theoretical
yield
Grain refinement, solid solution and precipitation strengthening mechanisms were reviewed as the contributing factors to the 0.2% yield strength of the present alloy. In the above discussion, analysis was done to show the contribution of each strengthening mechanism on the yield strength: Increment of yield strength associated with grain refinement was evaluated as 21 MPa mm − 1/2. The degree of solid solution strengthening was 35 – 50 MPa. Among the four mechanisms associated with strengthening by d% precipitates, the effect of the antiphase boundary by dislocation pairs is greatest. Moreover, the strength increment by S%precipitates was most suitably explained by the Orowan’s by-pass mechanism. The contribution of S% precipitates to the strength increment is found to be greater than that by d% precipitates. Some good correlations are observed between the shape parameters on S% precipitates and precipitation strengthening as shown in Fig. 8. A microstructure composed of densely distributed fine S% precipitates is desirable in improving the yield strength. As shown in Eq. (1), contributing factors to 0.2% yield strength of the present alloy can be added. Fig. 9 shows the comparison of measured and theoretically calculated yield strength. The measured strength and calculated one will be in better agreement if more accurate increments in strength by solid solution strengthening and M values are measured.
5. Conclusion Effects of thermomechanical processing, TMP, on the strengthening of the 2091 Al – Li alloy were investigated. The following results were obtained. The greater warm rolling reduction ratio and lower warm rolling temperature give a better tensile properties. All the tensile properties except elastic modulus increase as the cold rolling reduction ratio increases. When the cold rolling reduction ratio increases, trends of 0.2% yield strength, tensile strength and failure strain in the successive processing of annealing .
23
and cold rolling are the same as those of the cold rolling without annealing treatment. The annealing treatment before cold rolling lowers the value of tensile properties. Precipitation of d% and S% precipitates is found to be the main contribution to the strengthening of this Al–Li alloy: The creation of antiphase of d% precipitate by dislocation pairs and closely distributed fine S% precipitates are effective in improving the yield strength. The increment in strength by S% precipitation can be explained quantitatively by the Orowan’s mechanism.
Acknowledgements The authors would like to thank Alithium Co., Ltd. for supplying the testing materials and financial support.
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[20] J.C. Hung, A.J. Ardell, Proceedings of the 4th Aluminum – Lithium Conference, Paris, France, June, Les editions de Physique, 1987, pp. 373–383. [21] Y. Miura, A. Matsui, M. Furukawa, M. Nemoto, in: C. Baker, P.J. Gregson, S.J. Harris, C.J. Peel (Eds.), Aluminum – Lithium Alloys III, Institute of Metals, London, UK, 1986, pp. 427 – 434.
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The role of microstructures on the strengthening mechanisms of a thermomechanically processed 2091 Al–Li alloy Heon-Joo Kim a,*, Mitsuo Niinomi b a
Department of Metallurgical Engineering, Pukyong National Uni6ersity, San 100 Yongdang-dong, Nam-ku, Pusan 608 -739, South Korea b Department of Production System Engineering, Toyohashi Uni6ersity of Technology, Tempaku, Toyohashi 441, Japan Received 27 October 1999; received in revised form 9 February 2000
Abstract The relationship between tensile properties and microstructural parameters are discussed. Grain size, degree of recrystallization, size and distribution of precipitates were varied by thermomechanical processing. Tensile tests were performed and the microstructures were characterized using transmission electron microscopy (TEM). Quantitative analysis was carried out to explain the yield strength in connection with the change in the microstructural parameters. The primary strengthening precipitates of the under aged alloy were d% (Al3Li) and S% (Al2CuMg). The greater the warm rolling reduction ratio and the lower warm rolling temperature give a higher strength. Closely distributed fine S% precipitates were effective in improving the yield strength. The increment in yield strength by S% precipitation can be explained quantitatively by the Orowan’s mechanism. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Tensile properties of 2091 Al–Li alloy; Thermomechanical processing (TMP); d% (Al3Li) and S% (Al2CuMg) precipitates; Quantitative analysis of microstructural parameters; Orowan’s mechanism
1. Introduction Much attention has been paid to develop aluminum– lithium alloys that have the combination of lower density with increased stiffness and strength. The reduction in density and increase in stiffness are attributed to the lithium addition to aluminum which is a desirable property for aircraft structural applications. Consequently, numerous studies have been carried out on various mechanical and microstructural characteristics of these alloys [1 – 5]. A thermomechanical processing (TMP) [6] is a useful step in the grain refinement of Al – Li alloys [7]. Fine grains which are achieved by TMP, are effective in reducing the local stress concentration of coarse planar slips at grain boundaries by shortening the dislocation pile-up length in Al – Li alloys [8]. Depending on the variables in TMP, the degree of recrystallization as well as grain size can be controlled; * Corresponding author. Tel.: +82-51-6201480; fax: + 82-516250718. E-mail address: [email protected] (H.-J. Kim)
microstructures of completely recrystallized, partially recrystallized or completely unrecrystallized grains with different grain sizes can be obtained by the proper selection of temperature and deformation mode in TMP. The deficiencies in the Al–Li alloys can thus be improved through the use of microstructural control using TMP techniques. The mechanical properties of Al–Li alloys vary with the degree of recrystallization. According to the papers by Starke et al. [7] and Jata et al. [9], the unrecrystallized microstructure of 2020 Al– Li alloy increases the strain to failure compared with the recrystallized microstructure. The objective of this study is to analyze the effect of microstructural features on the strength of 2091 Al–Li alloy. A quantitative study was carried out, especially focusing on explaining the change in yield strength due to d% (Al3Li) and S% (Al2CuMg) precipitates strengthening mechanism. Changes of the microstructure in this alloy, i.e. grain size, degree of recrystallization and size, spacing, vol.% of precipitates were accomplished by changes in the TMP.
0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 0 ) 0 0 8 0 2 - 9
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2. Experimental procedures The material used in this study is 2091 Al –Li alloy, whose chemical composition in wt.% is 2.15Li, 2.10Cu, 1.49Mg, 0.12Zr and balance of Al. The iron, silicon and sodium are 0.06 wt.%, 0.03 wt.% and 3 ppm, respectively. The original ingots (1300×350 × 40 mm) of alloy were homogenized, then 5 mm planed off from both surfaces, processed with different TMPs and rolled to a final thickness of 6 mm plates. The alloy was processed with various TMPs as shown in Fig. 1 schematically. Detailed processes of the TMP are as follows. 1. The warm rolling reduction ratio was varied at 30, 45, 60 and 75% for a constant warm rolling temperature of 623 K. 2. The warm rolling temperature was varied to 573, 623 and 723 K for a constant warm rolling reduction ratio of 45%. 3. The cold rolling reduction ratio prior to solution treatment was varied at 0, 20 and 30%, after warm rolling at the condition of 623 K, 45%. 4. Cold rolling reduction ratio was varied at 0, 20
15
and 40%, after a constant 45% warm rolling condition at 623 K and annealing treatments, 673 K, 2 h. The maximum reduction ratio of cold rolling was 30% in case (c), compared with 40% in case of (d). This was due to the occurrence of edge cracking when an annealing step was not used before cold rolling. The TMP processes were designed to control the microstructural parameter of the precipitates, grain size, subgrain size and degree of recrystallization. Also 3% cold rolling just before artificial aging was conducted to promote an even distribution of the fine S% precipitates within all of the specimens. The underaging treatment, in which toughness of the alloy is increased by lesser precipitation at the grain boundaries [10], was finally carried out in the present TMP. The purpose of the design of these TMP processes can be found [11]. Tensile specimens were machined to plate type with the dimension of 2 mm thickness, 4 mm gage width and 20 mm gage length, with the axis parallel to the rolling direction of the plates. Tensile tests were performed at room temperature at a crosshead speed of 8.3× 10 − 6 m s − 1 using an Instron type machine. The 0.2% yield strength, ultimate tensile strength, Young’s modulus and fracture strain were measured from the tensile test results. Microstructures of the alloys were characterized using an optical microscopy and transmission electron microscopy, TEM. Polished samples for optical microscope observation were etched by immersion in Keller’s reagent. Thin foil for TEM were prepared with a twin-jet Tenupol apparatus using a 20% methyl acetate solution at 253 K and a potential difference of 20 V. Thickness of the TEM thin foils was measured according to the contamination spot method [12] for minimizing the error in measuring the microstructural parameters concerning precipitates. A computer image analyzer was also used to measure the shape and distribution of microstructural parameters.
3. Results
3.1. Microstructural obser6ation
Fig. 1. Schematic drawing of thermomechanical processings applied in this study.
Table 1 shows the observation results of representative samples by an optical microscope. The microstructures of all the samples treated with different warm rolling reduction ratios were completely unrecrystallized after aging. The grains became more fine as the warm rolling temperature decreased. But elongated grains were observed at higher temperatures.
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Table 1 Variation of degree of recrystallization and grain size with processing condition Designation
(a) Warm rolling reduction ratio change (At constant temperature, 623 K) (b) Warm rolling temperature change (At constant ratio, 45%) (c) Cold rolling reduction ratio change (At constant warm rolling condition, 623 K, 45%) (d) Annealing and cold rolling (At constant annealing condition, 673 K, 2 h)
Degree of recrystallization (%)
30% 75% 573 K 723 K 20% 30% 20% 40%
0 0 0 0 0 7.5 0 5.1
The grains were slightly coarser when the cold rolling reduction ratio was 20% in comparison with the 30% ratio. Owing to the effect of severe cold rolling (30%), some grains were recrystallized at the grain boundaries after successive processing. When annealing was carried out before cold rolling, some precipitates formed at the grain boundaries. Change in grain size was appreciable with cold rolling and no annealing. The grains were more refined in the 20% ratio sample than in the case of 40%, as compared to those with cold rolling without annealing. Some recrystallized grains also nucleated at the grain boundaries when the ratio was 40%. d% (Al3Li) and S% (Al2CuMg) precipitates which are primary strengthening precipitates, were confirmed by TEM analysis in the under aged 2091 Al – Li alloy. While this composite d% precipitates were present in some TMP conditions, needle or cylindrical S% precipitates were found under all conditions.
Grain size Longtudinal (mm)
Transverse (mm)
Short transverse (mm)
51.6 45.2 38.1 65.1 50.0 50.0 35.1 41.3
42.9 22.6 30.2 31.7 28.6 26.2 14.9 22.2
54.0 16.7 20.2 23.8 23.8 19.0 15.1 21.5
3.2. Effect of warm rolling reduction ratio on tensile property Fig. 2 exhibits the TEM micrographs of subgrains when the warm rolling reduction ratios are 30 and 75% for warm rolling temperature of 623 K, in which spherical d% and needle S% are precipitated. The microstructural parameters measured on TEM micrographs for each condition are listed in Table 2. Tensile properties for different warm rolling reduction ratios are shown in Table 3. While the change of 0.2% yield strength is negligible, tensile strength and fracture strain are increased as the warm rolling reduction ratio increases. A prominent increase in tensile strength and fracture strain is shown when the ratio is changed to 60%. At this condition, drastic changes followed, i.e. decrease in the diameter and average spacing between the nearest neighbor in d% precipitates and increase in the length and average spacing between the nearest neighbor in S% precipitates.
Fig. 2. Transmission electron microscopy (TEM) showing subgrain structure, d% and S% precipitates when warm rolling reduction ratio changed.
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Table 2 Effect of warm rolling reduction ratio on subgrain size and microstructural parameters Warm rolling reduction ratio (%)
Subgrain size (mm)
d% phase Volume fraction (%)
30 45 60 75
2.5 2.2 1.9 1.8
0.70 0.73 0.70 0.67
Diameter (nm) 32 31 27 26
Average spacing (nm) 77.7 79.0 64.4 59.1
S% phase
30 45 60 75
2.5 2.2 1.9 1.8
Volume fraction (%)
Length (nm)
0.23 0.22 0.20 0.20
70.5 83.2 98.5 111
Width (nm)
Average spacing (nm)
6 5 5 5
67.7 54.6 91.4 97.4
Table 3 Effect of warm rolling reduction ratio on tensile properties Warm rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
30 45 60 75
388 393 394 396
402 410 451 490
74.5 75.4 76.7 78.6
0.097 0.102 0.129 0.132
a
Yield strength is the stress at 0.2% offset strain.
Table 4 Effect of warm rolling temperature on subgrain size and microstructural parameters Warm rolling temperature (K)
Subgrain size (mm)
d% phase Volume fraction (%)
723 623 573
3.1 2.2 2.1
0.75 0.73 1.44
Diameter (nm) 33 31 15
Average spacing (nm) 67.9 79.0 30.5
S% phase
723 623 573
3.1 2.2 2.1
Volume fraction (%)
Length (nm)
0.22 0.22 0.15
79.9 83.2 125
Width (nm) 6 5 5
Average spacing (nm) 55.8 54.6 38.5
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Table 5 Effect of warm rolling temperature on tensile properties Warm rolling temperature (K)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
723 623 573
359 388 394
401 409 429
75.9 75.6 76.1
0.037 0.052 0.062
a
Yield strength is the stress at 0.2% offset strain.
Fig. 3. Transmission electron microscopy (TEM) micrographs showing subgrain structure, d% and S% precipitates when cold rolling reduction ratio changed.
Table 6 Effect of cold rolling reduction ratio before solution treatment on subgrain size and microstructural parameters Cold rolling reduction ratio (%)
0 20 30
Subgrain size (mm)
2.2 1.7 1.9
d% phase Volume fraction (%)
Diameter (nm)
Average spacing (nm)
0.73 1.38 1.71
31 19 22
79.0 65.9 55.7
Volume fraction (%)
Length (nm)
Width (nm)
0.22 0.11 0.16
83.2 60.1 57.8
S% phase
0 20 30
2.2 1.7 1.9
5 5 5
Average spacing (nm) 54.6 79.6 63.4
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smallest subgrain size and the most refined d% precipitates.
3.3. Effect of warm rolling temperature on tensile property Microstructural parameters for warm rolling temperature of 573, 623 and 723 K are shown in Table 4 when the warm rolling reduction ratio is kept constant at 45%. Table 5 shows the variation of the tensile properties as a function of the warm rolling temperature. Most of the tensile properties increase as the temperature de-creases. At a constant warm rolling reduction ratio, strength is maximum at the condition of the
3.4. Effect of cold rolling reduction ratio on tensile property Fig. 3 and Table 6 show TEM micrographs and microstructural parameters as the cold rolling reduction ratio changes, respectively. The effect of cold rolling ratio before solution heat treatment on tensile properties is shown in Table 7. All
Table 7 Effect of cold rolling reduction ratio before solution treatment on tensile properties Cold rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
0 20 30
389 394 402
419 447 488
76.6 75.6 75.0
0.049 0.062 0.067
a
Yield strength is the stress at 0.2% offset strain.
Table 8 Effect of cold rolling reduction ratio after annealing on subgrain size and microstructural parameters Cold rolling reduction ratio (%)
0 20 40
Subgrain size (mm)
2.3 1.8 2.1
d% phase Volume fraction (%)
Diameter (nm)
Average spacing (nm)
0.98 0.96 0.70
24 21 27
38.3 43.8 51.9
Volume fraction (%)
Length (nm)
Width (nm)
0.12 0.11 0.20
59.6 75.0 98.5
S% phase
0 20 40
2.3 1.8 2.1
5 5 5
Average spacing (nm) 63.8 101.2 55.3
Table 9 Effect of cold rolling reduction ratio after annealing on tensile properties Cold rolling reduction ratio (%)
Yield strengtha (MPa)
Ultimate tensile strength (MPa)
Young’s modulus (GPa)
Fracture strain
0 20 40
363 387 417
398 449 485
75.7 76.4 78.1
0.028 0.039 0.046
a
Yield strength is the stress at 0.2% offset strain.
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The resolved shear yield stress, ty, of the present alloy can be described as; ty = tAl + DtG.B. + DtS.S. + DtP
Fig. 4. Relationship between subgrain diameter, d − 1/2, and 0.2% yeild strength, s0.2, of pure aluminum.
(1)
where tAl is the resolved shear yield stress of the Al matrix, and tG.B., tS.S. and tP are the increments of the resolved shear stress associated with grain refinement, solid solution and precipitation strengthenings, respectively. In the following discussion, the contribution of each strengthening mechanism to the 0.2% yield strength of the present alloy is determined based on theoretical analysis or equations. Finally, each constituent strength is added up to compare with the experimental overall yield strength of the 2091 alloy.
4.1. Grain refinement strengthening Yield strength of the alloy related to grain size can be described by the following Hall–Petch equation [13];
Fig. 5. Schematic drawing of calculating solid solution strengthening.
the tensile properties except elastic modulus increase as the cold rolling reduction ratio increases.
3.5. Successi6e effect of annealing and cold rolling on tensile property Microstructural parameters for annealing and different cold rolling reduction ratios are shown in Table 8. Some precipitates appeared at grain boundaries though precipitates free zones were not detected. Table 9 shows the variation of the mechanical properties as a function of cold rolling reduction ratio, after annealing treatments. Tensile strength shows the maximum value at the most severe cold reduction ratio, 40%, that is, at the condition of the maximum volume fraction and the longest length of S% precipitates. All the values are, however, smaller than those without annealing treatments. These low values are due to the appearance of precipitates at grain boundaries.
4. Discussion The strengthening mechanism in the present alloy will be theoretically discussed in the following sections.
1 s0.2 = s0 + k (2)
d where s0 is the friction stress, k is the constant and d is the subgrain diameter. Fig. 4 shows the 0.2% yield strength of pure aluminum against d − 1/2. A linear relation is observed between s0.2 and d − 1/2 for pure aluminum. The data of pure Al reported by Wong et al. [14] are used in calculating the yield strength increment of the Al matrix in the following section. The slope, i.e. the increment of yield strength associated with grain refinement is 21 MPa mm − 1/2.
4.2. Solid solution strengthening The degree of solid solution strengthening for precipitation hardening aluminum alloys is, generally, evaluated when an alloy is in the as-quenched condition. The degree of solid solution strengthening was, therefore, modified by assuming that it decreased logarithmically with aging time as schematically shown in Fig. 5. The effect of solid solution strengthening could be assumed to weaken gradually and completely disappear by reaching at 24 h aging time because this alloy has exhibited peak hardness value at this aging time. The aging time adopted in this study was 10 h. The degree of solid solution strengthening for as-solutionized alloy was calculated by subtracting s0.2 and an increment in strength caused by grain refinement for pure aluminum which is the result of Section 4.1 from s0.2 of as-quenched alloy. The degree of non-modified solid solution strengthening of the alloy is from 140 to 150 MPa, while the modified one, i.e. the degree of solid solution strengthening is from 35 to 50 MPa. It was postulated that the rate of precipitation was constant
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during aging, and also it was the same in all the materials used.
(c) Strengthening due to misfit strain between matrix and d%;
4.3. Precipitation strengthening
Dto =
The main precipitates in the present alloy were found to be d% and S% from TEM observation. Strengthening mechanisms of Al – Li binary alloys by d% has been well defined, and good agreement between experimental and calculated values of strength has been reported [15]. The strengthening due to S% has not been examined in detail. The contributions of d% and S% to s0.2 of the present alloy will be discussed focussing on the strengthening mechanisms due to S%.
4.3.1. Strengthening by d% precipitates Dislocations move by shearing d%, because d% is coherent with the matrix. Noble et al. [15] have summarized that the strengthening of Al – Li type alloy by d% is associated with the following four mechanisms; (a) Strengthening due to the creation of antiphase boundary by dislocation pairs; g 4f −f (3) Dt0 = 2b p where Dt0 is the increment in resolved shear strength due to the creation of an antiphase boundary by dislocation pairs. (b) Strengthening due to the difference in shear modulus between the matrix and d%; DG 3DG r 3 · 0.8 − 0.143 ln · r · f DtG = 2 · 4p Gb b (4) where DtG is the increment in resolved shear strength due to the difference in shear modulus between matrix and d%, and DG is the difference in shear modulus between matrix and d%.
'
'
'!
"
(3G o 3 r f)
b
(5)
where Dto is the increment in resolved shear strength caused by misfit strain, G is the shear modulus of matrix (30 GPa), r is the effective radius of d%, o is the misfit strain (0.7 × 10 − 3) due to the precipitation of d%, f is the volume fraction of the d% precipitates and b is the Burgers’ vector (0.2864 nm). (d) Strengthening due to shearing of d% by dislocations; DtS =
6 f · gS · p r
(6)
where DtS is the increment in resolved shear strength due to shearing of d% by dislocations and gS is the surface energy of the matrix (0.2 J m − 2). The equation of s= Mt where M is the Taylor factor was used for converting resolved shear stress, t to tensile stress, s. Although M is a function of the processing parameter due to texture [16–18], M= 3.07 which had been proposed by Decker [13] and generally used for f.c.c. metals. The calculated values of the increment in tensile stress, according to the above-mentioned four strengthening mechanisms, are shown in Fig. 6. The increment in strength due to the creation of antiphase boundaries by dislocation pairs is greatest, approaching 75–107 MPa while those due to the misfit strain between d% precipitates and matrix, and shearing of d% precipitates by dislocations are very small.
4.3.2. Strengthening by S% precipitates The strengthening mechanism of commercial Al–Li alloys by S% precipitates has remained as unclear. The chemical strengthening due to the shearing of S% precipitates by dislocations and the strengthening due to the dislocation movement according to the Orowan’s by-pass mechanism were applied for explaining the strengthening of the present alloy by S% precipitates. 4.3.2.1. Chemical strengthening. The increments in resolved shear stress caused by weak and strong precipitates, Dtch,w, and Dtch,s are given by the following equations, respectively[19]; For weak precipitate, Dtch,w =
Fig. 6. Strengthening increments associated with d% precipitation obtained from various strengthening mechanisms.
'
4g 3bf GZD
(7)
where gi is the surface energy of precipitates, G (= 0.5Gb 2) is the line tension of dislocation, f is the
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stress by screw dislocation is given by the following Orowan’s equation modified by Ashby [22];
Z r0 DtSC = 2pD(1− y) Gb ln
Fig. 7. Comparison between experimental S% phase precipitation strengthening and theoretical one calculated by Orowan’s mechanism on yield strength.
Fig. 8. Relationship between precipitation strengthening increment for yield strength and S% phase parameters.
(9)
where D is the effective distance between S% precipitates, y is the Poisson’s ratio (0.33) and r0 (= 2b) is the radius of the dislocation core. The values of 1.41 times of the measured width and distance of S% precipitates were substituted for Z and D, respectively, because S% precipitated on (210)100 which crossed the matrix slip plane, (111), by 45° [23]. The measured increment in strength caused by S% precipitates can be given by subtracting the theoretically calculated strength increments caused by grain refinement, solid solution and d% precipitates from measured s0.2 of the aged alloy. Both modified and nonmodified strength increments by solid solution strengthening were taken into account in evaluating the measured increment in strength by S% precipitates. The calculated strength increment by S% precipitates which is based on the chemical strengthening mechanism of weak and strong precipitates and the measured one are compared. The calculated values are much smaller than the measured ones. The strength increment by S% precipitates, therefore, can not be explained by the chemical strengthening mechanism. The calculated strength increments by S% precipitates according to the Orowan’s mechanism are compared with measured ones. The calculated strength increment with linear correction of solid solution strengthening and measured one are shown in Fig. 7, in which Taylor factor M=3.07 was used. The calculated values are in good agreement with the measured values.
volume fraction of the S% precipitates, Z is the effective width of S% precipitates and D is the length of S% precipitates. For strong precipitate, Dtch,s = 1.62gi
'
f ZD
(8)
The value of gi was measured to be 1.18 J m − 2 using the reversion technique reported by Hung et al. [20] in the present study.
4.3.2.2. Strengthening due to Orowan’s mechanism. According to the report by Miura et al. [21], screw dislocations are predominant in the plastic deformation of Al–Li type alloys. The increment in resolved shear
Fig. 9. Comparison between measured and theoretical yield strength.
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Therefore, the increment in strength caused by S% precipitation can be explained by Orowan’s mechanism based on the above theoretical analysis. The increment of yield strength associated with S% precipitation by Orowan’s mechanism is assumed to be from 100–200 MPa.
4.4. Analysis strengths
of
measured
and
theoretical
yield
Grain refinement, solid solution and precipitation strengthening mechanisms were reviewed as the contributing factors to the 0.2% yield strength of the present alloy. In the above discussion, analysis was done to show the contribution of each strengthening mechanism on the yield strength: Increment of yield strength associated with grain refinement was evaluated as 21 MPa mm − 1/2. The degree of solid solution strengthening was 35 – 50 MPa. Among the four mechanisms associated with strengthening by d% precipitates, the effect of the antiphase boundary by dislocation pairs is greatest. Moreover, the strength increment by S%precipitates was most suitably explained by the Orowan’s by-pass mechanism. The contribution of S% precipitates to the strength increment is found to be greater than that by d% precipitates. Some good correlations are observed between the shape parameters on S% precipitates and precipitation strengthening as shown in Fig. 8. A microstructure composed of densely distributed fine S% precipitates is desirable in improving the yield strength. As shown in Eq. (1), contributing factors to 0.2% yield strength of the present alloy can be added. Fig. 9 shows the comparison of measured and theoretically calculated yield strength. The measured strength and calculated one will be in better agreement if more accurate increments in strength by solid solution strengthening and M values are measured.
5. Conclusion Effects of thermomechanical processing, TMP, on the strengthening of the 2091 Al – Li alloy were investigated. The following results were obtained. The greater warm rolling reduction ratio and lower warm rolling temperature give a better tensile properties. All the tensile properties except elastic modulus increase as the cold rolling reduction ratio increases. When the cold rolling reduction ratio increases, trends of 0.2% yield strength, tensile strength and failure strain in the successive processing of annealing .
23
and cold rolling are the same as those of the cold rolling without annealing treatment. The annealing treatment before cold rolling lowers the value of tensile properties. Precipitation of d% and S% precipitates is found to be the main contribution to the strengthening of this Al–Li alloy: The creation of antiphase of d% precipitate by dislocation pairs and closely distributed fine S% precipitates are effective in improving the yield strength. The increment in strength by S% precipitation can be explained quantitatively by the Orowan’s mechanism.
Acknowledgements The authors would like to thank Alithium Co., Ltd. for supplying the testing materials and financial support.
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