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Investigation of the quenching sensitivity of forged 2A14 aluminum alloy by time-temperature-tensile properties diagrams
Journal of Alloys and Compounds 728 (2017) 1239e1247
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Investigation of the quenching sensitivity of forged 2A14 aluminum alloy by time-temperature-tensile properties diagrams Yuxun Zhang a, b, Youping Yi a, b, *, Shiquan Huang a, b, Fei Dong a, b, Huimin Wang a, b a b
State Key Laboratory of High Performance Complex Manufacturing, Central South University, Changsha, 410083, China School of Mechanical and Electrical Engineering, Central South University, Changsha, 410083, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 January 2017 Received in revised form 2 September 2017 Accepted 4 September 2017 Available online 7 September 2017
To guide the design of quenching technologies for forged 2A14 aluminum alloys, the quenching sensitivity of this material was investigated with time-temperature-tensile properties (TTP) diagrams. The TTP diagrams were determined by interrupted quenching experiments followed by tensile testing. The microstructures were observed with TEM and EDX. The results indicate that the quenching sensitivity temperature range is from 300 to 390 C and the nose temperature is about 350 C. The tensile strength and yield strength decrease rapidly in the quenching sensitivity temperature range, and they decrease most rapidly at the nose temperature. In addition, the yield strengths decrease faster than the tensile strengths with increasing isothermal holding times. At the nose temperature of 350 C, quench-induced rod-shaped q0 particles precipitate and grow rapidly because of the high diffusion rate of the solute and the high phase-nucleation rate. The amount and size of the rod-shaped q0 particles increase when the holding time is prolonged. This increase results in a decrease of supersaturated solid solution after the quenching, which reduces the density of the needle-shaped fine age-induced precipitates q0 and q00 (strengthening particles) and increase the breadth of precipitate free zone (PFZ) after the aging treatments. © 2017 Published by Elsevier B.V.
Keywords: 2A14 aluminum alloy Quenching Tensile properties Microstructure Phase transitions
1. Introduction Heat-treatable aluminum alloys are widely used to fabricate forged components used in the aerospace and aircraft industry for weight reduction. Solution heat, quenching, and aging treatments are applied to obtain desirable mechanical properties [1]. In order to improve the mechanical properties, rapid cooling rates are necessary to minimize quench-induced precipitation [2e4]. However, rapid cooling rates produce excessive residual stress and distortion, which degrades the mechanical properties and dimensional stability [5]. Cooling rates are needed to decline to reduce the residual stress [6]. Fortunately, phases of materials usually precipitate rapidly in the quenching sensitivity temperature range, and precipitation will be slower in other temperature ranges compared to the critical range [7e9]. Consequently, the mechanical properties and residual stress can be balanced by choosing a suitable cooling profile. For example, a cooling profile may incorporate quenching at
* Corresponding author. State Key Laboratory of High Performance Complex Manufacturing, Central South University, Changsha, 410083, China. E-mail addresses: [email protected] (Y. Zhang), [email protected] (Y. Yi). http://dx.doi.org/10.1016/j.jallcom.2017.09.041 0925-8388/© 2017 Published by Elsevier B.V.
critical cooling rates within the quenching sensitivity temperature range, along with slow cooling rates in the other temperature ranges. The quenching sensitivity temperature range and critical cooling rates might vary for different materials. It is important to study the quenching sensitivity of the material in order to optimize the quenching technology. Time-Temperature-Transformation/Properties (TTT/TTP) diagrams are usually applied to investigate quenching sensitivity [10e13] and are often acquired through interrupted quenching experiments. For example, Li et al. [14] used these diagrams to investigate the quenching sensitivity of a 6351 aluminum alloy and showed that its quenching sensitivity range is from 230 to 430 C. With the use of TTP curves, Shang et al. [11] found that the quenching sensitivity temperature range of a 6082 aluminum alloy is from 250 to 440 C with the nose temperature equal to 360 C. Hardness is an important mechanical property, and most researchers measure it to determine the TTP curves to investigate the evolution of the mechanical properties during quenching. However, tensile strength rather than hardness is the reason that aluminum alloys are chosen for mechanical applications. There is relatively little information in the published literature on the evolution of
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tensile properties and microstructures of the 2A14 aluminum alloy during quenching. In this work, tensile properties and microstructural evolution in a forged 2A14 aluminum alloy during quenching were investigated by performing interrupted quenching experiments. The tensile properties were measured to investigate the evolution of the mechanical properties during quenching. Based on the results, the TTP diagrams were drawn and the quenching sensitivity of the material was studied. In addition, the microstructural evolution during the isothermal treatments was observed with transmission electron microscopy (TEM), energy-dispersive X-ray spectroscopy (EDX), and X-ray diffraction (XRD). 2. Materials and methods 2.1. Interrupted quenching experiments Two kinds of 2A14 aluminum alloy specimens were used in this work. Small samples (15 15 4 mm) were used to measure the electrical conductivity, and larger samples (2 mm 8 mm with a gage length of 30 mm) were used to measure the tensile properties. The specimens were machined from a commercially forged component along the rolling direction. The chemical composition of the samples is listed in Table 1. After solution treatment in an air furnace at 500 C for 3 h, the specimens were transferred within 3 s to a salt bath at a designated temperature from 250 to 480 C. After being held isothermally for the required time (5e1200 s), the specimens were immediately quenched in room temperature water (about 20 C). Subsequent artificial aging treatments were carried out at 165 C for 9 h within 4 h after quenching. The electrical conductivities were measured by a D60K within 1.5e3 h after the quenching treatments. The tensile properties were measured in accordance with the GB/T 1685-2013 standard after the artificial aging treatments. Three tensile test pieces were used for every set of experimental conditions. The samples were tested at a strain rate of 0.0011 s1. 2.2. Microstructural observations The microstructures of the samples were observed using a Titan G2 60-300 transmission electron microscope. An EDX system attached to the TEM was employed in selected domains. Thin foils for TEM were prepared by mechanically polishing the samples to thicknesses of 150 mm followed by double-jet electro-chemical polishing in a solution of 30% nitric acid and 70% methanol at 20 C. The Al-Y3800 was used to perform the XRD analysis of the samples. 3. Results and analysis
without isothermal holding treatment was taken as the minimum value (gmin). The value of electrical conductivity [49% (IACS)] of the sample held isothermally for 48 h at 350 C was taken as the maximum value (gmax).
f ¼
g gmin gmax gmin
(1)
Fig. 1 shows the electrical conductivities and phase transformation volume fractions of the as-quenched samples. The samples were held isothermally at different temperatures for various times. As shown in Fig. 1 (c), the phase transformation volume fraction increases with increasing isothermal holding time. It increases sharply at the initial isothermal stage, and then the change rate decreases with additional isothermal holds. Finally, it remains almost stable after a long isothermal duration. In the case of isothermal holding at 350 C, the first 60 s are required to transform 33% of the phase volume fraction; then an additional 240 s (from 60 s to 300 s) are needed to transform another 31% phase volume fraction. Finally, the value remains almost constant at 72% after isothermal holding for 600 s. In addition, the transformation rate depends on temperature; it increases with increasing temperature below 350 C, and decreases with increasing temperature above 350 C. With isothermal holding for 300 s, the phase transformation volume fraction is 64% at 350 C, while at 250 C and 470 C, the values are 25% and 9%, respectively. The tensile strengths and yield strengths of the as-aged samples were measured to investigate the evolution of the mechanical properties during quenching. Fig. 2 shows the tensile strengths of the samples with an isothermal holding treatment at different temperatures for different times. The results indicate that the tensile strength decreases rapidly with prolonged isothermal duration at the initial stage, and then the rate of change decreases as the duration increases. At a moderate temperature of 350 C, it takes 30 s to decrease from 480 MPa to 391 MPa, and then it takes 240 s to decrease from 391 MPa to 278 MPa. Furthermore, the rate of change of tensile strength also depends on temperature. Below 350 C, the decreasing rates increase with decreasing temperatures; conversely, the decreasing rate decreases with increasing temperatures from 350 to 480 C. When time is held constant at 300 s and tensile strength is analyzed at 250, 350, and 480 C, it is apparent that the rate of change of the tensile strength is not affected as much at low and high temperatures as it is at the moderate temperature. At 250 C, the initial tensile strength of 480 MPa only decreased to 441 MPa and at 480 C it only decreased to 467 MPa, whereas at the moderate temperature of 350 C, the tensile strength decreased to 278 MPa after 300 s. Fig. 3 shows the yield strengths of the as-aged samples with an isothermal holding treatment at different temperatures for different times. The trend of the yield strengths with temperature and isothermal time is similar to that of the tensile strengths.
3.1. Phase transformation fractions and tensile properties In this work, the time-temperature-transformation fraction diagram was used to show the evolution of the phase transformation during the isothermal holding treatment. The values of the electrical conductivities (g) were transformed to phase transformation volume fractions (f) by using Eq. (1) [15]. In this equation, the value of electrical conductivity [34% (IACS)] of the as-quenched sample
Table 1 Chemical composition of the studied material (wt%). Cu
Mg
Mn
Si
Fe
Ni
Zn
Ti
Al
4.28
0.6
0.81
0.94
0.15
0.003
0.01
0.04
Balance
3.2. TTP diagrams The average tensile strength and yield strength of the test samples without any isothermal holding time were 480 MPa and 430 MPa, respectively. These values were taken as the maximum values, and then the contours of percentage of the tensile properties for each sample as a function of different isothermal treatments were constructed. Fig. 4 shows the TTP contours for the tensile strengths and yield strengths. The results indicate that the shape of the contours is a “C” type. The tensile strengths and yield strengths decrease most rapidly at 350 C. At a given temperature, the breadth of the iso-contours from 100% to 60% of the yield strengths is smaller than that of the tensile strengths. This means that the
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Fig. 1. Electrical conductivities and phase transformation volume fractions of the as-quenched samples with isothermal holding treatment at different temperatures from 250470 C for different times: (a, b) electrical conductivities; (b) phase transformation volume fractions.
Fig. 2. Tensile strengths of the as-aged samples with isothermal holding treatment at different temperatures for different times: (a) 480-350 C; (b) 350-250 C.
yield strengths decrease faster than the tensile strengths with increased isothermal holding times. TTP curves derived from the tensile strengths were used to further investigate the quenching sensitivity of the studied material. The TTP curve can be described by Eq. (2) [16]. By using this
equation, the critical time (tc) to produce a certain percentage of the material at a maximum tensile properties can be calculated, since the parameter tc is related to the amount of quench-induced phase transformation volume fraction.
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Fig. 3. Yield strengths of the as-aged samples with an isothermal holding treatment at different temperatures for different times: (a) 480-350 C; (b) 350-250 C.
Fig. 4. TTP contours for 2A14 aluminum alloys: (a) tensile strengths; (b) yield strengths.
1 tc ¼ k1 k2 exp RT
k3 k24 ðk4 TÞ2
!! þ k5
(2)
where k1 is the natural logarithm of the untransformed fractions; k2 is a constant related to the reciprocal of the number of nucleation sites (s); k3 is a constant related to the energy required to form a nucleus (J$mol1); k4 is a constant related to the solvus temperature; k5 is a constant related to the activation energy for diffusion (J$mol1); R is the gas constant (J$mol1$K1); and T is the absolute temperature (K). According to the tensile strength results, the times to achieve 95% of the maximum tensile strength at different temperatures were used to fit Eq. (2). The coefficients in Eq. (2) were evaluated by an iterative non-linear fitting method. The coefficients k2-k5 are shown in Table 2. By using Eq. (2) and the coefficients in Table 2, the TTP curves of the studied material were derived at 95%, 90%, 85%, and 80% of their maximum levels, as shown in Fig. 5. The nose Fig. 5. Time-temperature-tensile strengths (TTP) curves.
Table 2 Coefficients of Eq. (2). k3/J$mol1
k2/s 1.21 10
12
9.71 10
3
k4/K
k5/J$mol1
1034
108462
temperature is about 350 C. The tensile strength decreases rapidly in the temperature range from 300 to 390 C, while, it changes
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Fig. 6. TEM micrographs of the as-quenched samples without isothermal holding: (a) bright-field TEM; (b) EDX spectra at position A.
Fig. 7. TEM micrographs of the as-quenched samples after isothermal holding for different times at 350 C: (a) 30 s (dark field); (b) EDX spectra at position P1 in (a); (c) 1200 s (bright field); (d) XRD analyses after holding for different times at 350 C.
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Fig. 8. {001}Al TEM micrographs of the as-aged samples after isothermal holding for different times at 350 C: (a, b) 10 s; (b) the SAD pattern of (a); (d) 30 s; (e) 300 s.
slowly in other temperature ranges, especially at temperatures above 450 C. It requires more than 450 s for the tensile strength to decrease to 95% of the maximum value at 480 C, but less than 10 s to decrease to 95% of the maximum value at 350 C. To obtain 95% of the maximum tensile strength for the forged 2A14 aluminum alloy components, the critical average cooling rate was about 7 C s1 at temperatures from 300 to 390 C. 3.3. Microstructural observation In order to clarify the microscopic mechanism of quenching sensitivity, TEM, EDX, and XRD analyses were used to investigate the precipitation behavior of the studied material at 450 C, 350 C, and 250 C with different isothermal holding times. Fig. 6 shows the microstructure of the as-quenched sample without any isothermal holding time. As shown in Fig. 6 (a), there are a few rod-shaped dispersed particles (arrow A). Their lengths are in the range of 100e350 nm and their diameters are in the range of 50e120 nm. According to the EDX results [Fig. 6 (b)], the precipitates are enriched with Cu/Mn. Referring to Ref. [17], it can be concluded that this phase is the T (Al20Cu2Mn3) phase. Except for the T phase, there are no other precipitate phases in the matrix. Fig. 7 shows the microstructure of the as-quenched sample after isothermal holding at 350 C for 30 s and 1200 s. The size and density of the quench-induced precipitate phase increases with increasing isothermal holding time. In the case of the sample with
isothermal holding for 30 s, the T phase in Fig. 7 (a) shows no obvious coarsening, compared to the sample without isothermal treatment [Fig. 6 (a)]. However, a large number of quench-induced rod-shaped precipitates with a length in the range of 70e350 nm (arrow P1) appeared in the matrix. According to the EDX analysis [Fig. 7 (b)], the precipitates are enriched with Cu. Referring to Refs. [12,18], these should be q0 phases. In the case of the sample with isothermal holding for 1200 s, the density and size of the quench-induced rod-shaped phase increases dramatically, compared to the sample with isothermal holding for 30 s, as shown in Fig. 7 (c). The long isothermal holding time caused the sample to lose more solutes, which reduced the volume fraction and density of fine age-induced precipitates and retarded the growth of them in the subsequent aging process. This resulted in poorer mechanical properties, specifically a lower yield strength. The XRD result also shows in Fig. 7 (d) that with the increased isothermal holding times, the density of the quench-induced precipitated phases (Al2Cu) increased, which is consistent with the microstructural observations. Fig. 8 shows the microstructure of as-aged samples with different isothermal holding times at 350 C. In the case of the sample subjected to isothermal holding for 10 s, in addition to the coarse T-phases, another needle-shaped precipitate phase appeared with a length of about 10e100 nm, as shown in Fig. 8 (a). This new phase is distributed homogeneously along the {001}Al planes in the matrix. The corresponding SAD pattern shows that the
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Fig. 9. {001}Al TEM micrographs of the as-aged samples after isothermal holding for different times at 350 C: (a) 0 s; (b) 10 s; (c) 30 s; (d) EDS mapping images for the sample subjected to isothermal holding for 10 s.
streaks are discontinuous [Fig. 8(c)]. Based on prior literature [17,19e21], Fig. 8(a, c) suggests that the age-induced needle-shaped precipitate is a q0 fine precipitate. As shown in the HRTEM image [Fig. 8 (b)], there are some small needle-shaped precipitates along the {001}Al planes in the matrix, comprising a triple copper-layer structure with a length of about 15 nm. According to the literature [22], this should be a q00 fine precipitate, which means that the q0 and q00 precipitated phases existed together in the matrix. With the increase of isothermal times to 30 s, the density of small ageinduced needle-shaped fine q0 phases is obviously reduced and the particles are distributed unevenly in the matrix, as shown in Fig. 8 (d). In addition, the density of coarse rod-shaped q0 phases increases and there is a large precipitate free zone (PFZ) around the coarse phases. When the isothermal time is extended to 300 s, the size of the coarse precipitates (q phases) increases sharply to about 500 nm in length and the fine precipitates almost disappear. Fig. 9 shows precipitated phases along the grain boundaries of the as-aged samples with different isothermal holding times at 350 C. In the case of the sample without any isothermal holding treatment, the rod-shaped precipitates, with a length of about 100 nm, are distributed discontinuously along the grain boundary, and the PFZ zone is smaller than 40 nm. In the case of the sample that was subjected to isothermal holding for 10 s, the precipitated phases along the grain boundary are distributed continuously, and the PFZ zone increases to about 160 nm. For the sample with isothermal holding for 30 s, the precipitates grow to about 200 nm in length, and the PFZ zone increases to more than 240 nm. As shown in Fig. 9 (d), the shapes of the precipitated phases along the grain boundary are rectangular, and they are enriched with Cu. We inferred that the phase is q0 (q) (Al2Cu) in the grain boundaries. Fig. 10 shows the microstructures of the samples subjected to isothermal holding for 30 s at 450 C and at 250 C. In contrast to
the microstructures of the 350 C samples in Figs. 7 (a) and 10 (a) shows that there are no quench-induced rod-shaped precipitated phases, except for the coarse T phase, in the matrix of the asquenched sample isothermally held at 450 C. On the other hand, there are some quench-induced precipitated phases, again except for the coarse T phase, in the matrix of the as-quenched sample isothermally held at 250 C, as shown in Fig. 10 (b). However, the amount is smaller than at 350 C. As shown in Fig. 10 (cef), in the case of the as-aged samples isothermally treated at 450 and 250 C for 30 s, the needle-shaped q0 precipitated phases are smaller in length and distributed uniformly in the matrix, compared to the sample isothermally treated at 350 C for the same time [Figs. 8 (d) and 9 (c)]. In addition, the precipitated phases along the grain boundaries are distributed discontinuously, and the PFZ zone is smaller. The results indicate that the quench-induced phase precipitation rate is much smaller at 450 C and 250 C than at 350 C. Consequently, for the same isothermal holding time, the degree of supersaturation of the as-quenched samples treated at 450 C and 250 C remains high, which increase the density and volume fraction of fine age-induced precipitates after aging treatment, which improves the mechanical performance. This result is in agreement with the results of the mechanical properties tests, as shown in Figs. 2 and 3. 4. Discussion This work shows that for the studied material, increasing the isothermal holding times increases the phase transformation fractions during quenching, and reduces the tensile strengths and yield strengths after aging treatments. These material properties change rapidly during the quenching sensitivity temperature range from 300-390 C, and most rapidly at the nose temperature of 350 C.
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Fig. 10. {001}Al TEM bright field micrographs of the as-quenched and as-aged 2A14 aluminum alloy after isothermally holding for 30 s at 450 C and 250 C: (a) as-quenched sample isothermally held at 450 C; (b) as-quenched sample isothermally held at 250 C; (c, d) as-aged sample isothermally held at 450 C; (e, f) as-aged sample isothermally held at 250 C.
When the quench-induced phase transformation fractions increase, the density of fine age-induced precipitates and the breadth of PFZ zone decrease after the aging treatments. The metals and alloys can be strengthened by the fine secondphase particles, and the strengths increase with the density and size of fine second-phase particles [23,24]. For this studied material, the spaces between the quench-induced precipitates are big, and it is easy for the dislocations pass through it [25]. The fine ageinduced precipitates are the main strengthening particles. The density and volume fraction of the age-induced precipitates decrease with increasing quench-induced phase transformation fraction during quenching. In addition, the breadth of the PFZ zone increases with increasing quench-induced phase transformation fraction. Because the quench-induced phases are coarse and grow during the subsequent aging treatment, which consume the solute atoms in the matrix, especially the atoms around the coarse quench-induced precipitated phases. During isothermal quenching, the phase transformation fraction is determined by the nucleation and decomposition of the supersaturated solid solution. The quench sensitivity is determined by the concentration of alloy elements and nucleation sites [26,27]. The nucleation rate during quenching can be described by Eq. (3) [12,28]:
I ¼ cexpð DG* þ Q =ðRTÞ
(3)
where c is a constant, DG* is the nucleation activation energy, Q is the solute diffusion activation energy, R is the gas constant, and T is the absolute temperature. Eq. (3) shows that the phase-nucleation rate increases as the temperature is lowered. Meanwhile, the diffusion process significantly affects precipitation and growth of the quench-induced phases, and diffusion rate decreases with decreasing temperature. The precipitation processes during quenching are determined by temperature, and the effects are complicated. In general, at high temperatures near the solution temperature, the diffusion rates of solute atoms are high; however, the degree of supersaturation of solute atoms is very low and the nucleation rate is very small. Consequently, the precipitation rate is small. On the other hand, the nucleation rate is high at low temperatures because of the high degree of supersaturation, but the diffusion rate is slow, which results in a low precipitation rate. Meanwhile, at the intermediate temperatures, the precipitation rate is high, because both the nucleation rate and the diffusion rate are high, which explains why the quenching sensitivity temperature range occurs at intermediate temperatures. Even though the evolution of microstructures and tensile properties of the forged 2A14 aluminum alloy during quenching-
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aging treatment has been analyzed here, there are still some insufficiencies. First, this work did not explore why the yield strength decreases faster than tensile strength with increasing isothermal holding time. Second, the influence of heating treatments on the evolution of microstructures and mechanical properties was not investigated. For the studied material, the heating treatment varies for different aims [29,30]. In addition, heating treatments and the amount of plastic deformation affects the quenching sensitivity of aluminum alloys [31,32]. For example, T73 treatment decreases the nose temperature of the TTP curve by 40 C for the 7075 aluminum alloy, compared to T6 treatment [31]. In order to expand the application of this study, the influence of heating treatments on the microstructural evolution and mechanical properties should be investigated in future work. 5. Conclusion In this work, the quenching sensitivity of forged 2A14 aluminum alloy was investigated by the time-temperature-tensile strength/ yield strength diagrams. TEM, EDX, and EDS methods were applied to observe the changes in microstructure. Based on the results and analysis, several conclusions can be drawn: 1) The nose temperature of the TTP contours and curves of the forged 2A14 aluminum alloy is about 350 C, and the quenching sensitivity temperature zone is from about 300 to 390 C. The mechanical properties degrade rapidly in the quenching sensitivity temperature range and most rapidly at 350 C, but the changes became slow above or below this temperature. 2) The heterogeneous precipitation of coarse equilibrium q0 and q phases reduces the concentration of solutes, and they precipitate rapidly at temperatures from 300 to 390 C during isothermal holding treatments, especially at 350 C. The increasing quench-induced phase transformation fractions reduce the density and volume fraction of the fine age-induced needle-shaped q0 and q00 precipitates, and increase the size and the breadth of the PFZ zone, which results in a lower mechanical performance. 3) Yield strengths decrease faster with increasing isothermal holding time than tensile strengths. 4) The critical average cooling rate is about 7 C/s in the quenching sensitivity temperature range to obtain 95% of the maximum tensile strength. 5) In order to balance the residual stress and mechanical performances of the forged 2A14 aluminum alloy components, it is suggested that the cooling rate should be enhanced in the temperature range from 300 to 390 C and suppressed in the other temperature ranges. Acknowledgments This work was financially supported by State Key Laboratory of High Performance Complex Manufacturing (zzyjkt2014-02) and the fund of Jiangsu Province for the transformation of scientific and technological achievements (BA2015075). References [1] Y. Xiao, S. Xie, J. Liu, T. Wang, Practical Handbook of Aluminum Technology, Metallurgical industry Press, Beijing, 2005. [2] B. Milkereit, M.J. Starink, Quench sensitivity of Al-Mg-Si alloys: a model for linear cooling and strengthening, Mater. Des. 76 (2015) 117e129. [3] B.C. Shang, Z.M. Yin, G. Wang, B. Liu, Z.Q. Huang, Investigation of quench sensitivity and transformation kinetics during isothermal treatment in 6082
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Deng, Investigation of quench sensitivity of high strength AleZneMgeCu alloys by timeetemperatureproperties diagrams, Mater. Des. 31 (2010) 3116e3120. [32] L. Sheng-dan, L. Cheng-bo, O. Hui, D. Yun-lai, Z. Xin-ming, L. Xing-xing, Quench sensitivity of ultra-high strength 7000 series aluminum alloys, Chin. J. Nonferrous Metals 23 (2013) 927e938.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Investigation of the quenching sensitivity of forged 2A14 aluminum alloy by time-temperature-tensile properties diagrams Yuxun Zhang a, b, Youping Yi a, b, *, Shiquan Huang a, b, Fei Dong a, b, Huimin Wang a, b a b
State Key Laboratory of High Performance Complex Manufacturing, Central South University, Changsha, 410083, China School of Mechanical and Electrical Engineering, Central South University, Changsha, 410083, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 January 2017 Received in revised form 2 September 2017 Accepted 4 September 2017 Available online 7 September 2017
To guide the design of quenching technologies for forged 2A14 aluminum alloys, the quenching sensitivity of this material was investigated with time-temperature-tensile properties (TTP) diagrams. The TTP diagrams were determined by interrupted quenching experiments followed by tensile testing. The microstructures were observed with TEM and EDX. The results indicate that the quenching sensitivity temperature range is from 300 to 390 C and the nose temperature is about 350 C. The tensile strength and yield strength decrease rapidly in the quenching sensitivity temperature range, and they decrease most rapidly at the nose temperature. In addition, the yield strengths decrease faster than the tensile strengths with increasing isothermal holding times. At the nose temperature of 350 C, quench-induced rod-shaped q0 particles precipitate and grow rapidly because of the high diffusion rate of the solute and the high phase-nucleation rate. The amount and size of the rod-shaped q0 particles increase when the holding time is prolonged. This increase results in a decrease of supersaturated solid solution after the quenching, which reduces the density of the needle-shaped fine age-induced precipitates q0 and q00 (strengthening particles) and increase the breadth of precipitate free zone (PFZ) after the aging treatments. © 2017 Published by Elsevier B.V.
Keywords: 2A14 aluminum alloy Quenching Tensile properties Microstructure Phase transitions
1. Introduction Heat-treatable aluminum alloys are widely used to fabricate forged components used in the aerospace and aircraft industry for weight reduction. Solution heat, quenching, and aging treatments are applied to obtain desirable mechanical properties [1]. In order to improve the mechanical properties, rapid cooling rates are necessary to minimize quench-induced precipitation [2e4]. However, rapid cooling rates produce excessive residual stress and distortion, which degrades the mechanical properties and dimensional stability [5]. Cooling rates are needed to decline to reduce the residual stress [6]. Fortunately, phases of materials usually precipitate rapidly in the quenching sensitivity temperature range, and precipitation will be slower in other temperature ranges compared to the critical range [7e9]. Consequently, the mechanical properties and residual stress can be balanced by choosing a suitable cooling profile. For example, a cooling profile may incorporate quenching at
* Corresponding author. State Key Laboratory of High Performance Complex Manufacturing, Central South University, Changsha, 410083, China. E-mail addresses: [email protected] (Y. Zhang), [email protected] (Y. Yi). http://dx.doi.org/10.1016/j.jallcom.2017.09.041 0925-8388/© 2017 Published by Elsevier B.V.
critical cooling rates within the quenching sensitivity temperature range, along with slow cooling rates in the other temperature ranges. The quenching sensitivity temperature range and critical cooling rates might vary for different materials. It is important to study the quenching sensitivity of the material in order to optimize the quenching technology. Time-Temperature-Transformation/Properties (TTT/TTP) diagrams are usually applied to investigate quenching sensitivity [10e13] and are often acquired through interrupted quenching experiments. For example, Li et al. [14] used these diagrams to investigate the quenching sensitivity of a 6351 aluminum alloy and showed that its quenching sensitivity range is from 230 to 430 C. With the use of TTP curves, Shang et al. [11] found that the quenching sensitivity temperature range of a 6082 aluminum alloy is from 250 to 440 C with the nose temperature equal to 360 C. Hardness is an important mechanical property, and most researchers measure it to determine the TTP curves to investigate the evolution of the mechanical properties during quenching. However, tensile strength rather than hardness is the reason that aluminum alloys are chosen for mechanical applications. There is relatively little information in the published literature on the evolution of
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tensile properties and microstructures of the 2A14 aluminum alloy during quenching. In this work, tensile properties and microstructural evolution in a forged 2A14 aluminum alloy during quenching were investigated by performing interrupted quenching experiments. The tensile properties were measured to investigate the evolution of the mechanical properties during quenching. Based on the results, the TTP diagrams were drawn and the quenching sensitivity of the material was studied. In addition, the microstructural evolution during the isothermal treatments was observed with transmission electron microscopy (TEM), energy-dispersive X-ray spectroscopy (EDX), and X-ray diffraction (XRD). 2. Materials and methods 2.1. Interrupted quenching experiments Two kinds of 2A14 aluminum alloy specimens were used in this work. Small samples (15 15 4 mm) were used to measure the electrical conductivity, and larger samples (2 mm 8 mm with a gage length of 30 mm) were used to measure the tensile properties. The specimens were machined from a commercially forged component along the rolling direction. The chemical composition of the samples is listed in Table 1. After solution treatment in an air furnace at 500 C for 3 h, the specimens were transferred within 3 s to a salt bath at a designated temperature from 250 to 480 C. After being held isothermally for the required time (5e1200 s), the specimens were immediately quenched in room temperature water (about 20 C). Subsequent artificial aging treatments were carried out at 165 C for 9 h within 4 h after quenching. The electrical conductivities were measured by a D60K within 1.5e3 h after the quenching treatments. The tensile properties were measured in accordance with the GB/T 1685-2013 standard after the artificial aging treatments. Three tensile test pieces were used for every set of experimental conditions. The samples were tested at a strain rate of 0.0011 s1. 2.2. Microstructural observations The microstructures of the samples were observed using a Titan G2 60-300 transmission electron microscope. An EDX system attached to the TEM was employed in selected domains. Thin foils for TEM were prepared by mechanically polishing the samples to thicknesses of 150 mm followed by double-jet electro-chemical polishing in a solution of 30% nitric acid and 70% methanol at 20 C. The Al-Y3800 was used to perform the XRD analysis of the samples. 3. Results and analysis
without isothermal holding treatment was taken as the minimum value (gmin). The value of electrical conductivity [49% (IACS)] of the sample held isothermally for 48 h at 350 C was taken as the maximum value (gmax).
f ¼
g gmin gmax gmin
(1)
Fig. 1 shows the electrical conductivities and phase transformation volume fractions of the as-quenched samples. The samples were held isothermally at different temperatures for various times. As shown in Fig. 1 (c), the phase transformation volume fraction increases with increasing isothermal holding time. It increases sharply at the initial isothermal stage, and then the change rate decreases with additional isothermal holds. Finally, it remains almost stable after a long isothermal duration. In the case of isothermal holding at 350 C, the first 60 s are required to transform 33% of the phase volume fraction; then an additional 240 s (from 60 s to 300 s) are needed to transform another 31% phase volume fraction. Finally, the value remains almost constant at 72% after isothermal holding for 600 s. In addition, the transformation rate depends on temperature; it increases with increasing temperature below 350 C, and decreases with increasing temperature above 350 C. With isothermal holding for 300 s, the phase transformation volume fraction is 64% at 350 C, while at 250 C and 470 C, the values are 25% and 9%, respectively. The tensile strengths and yield strengths of the as-aged samples were measured to investigate the evolution of the mechanical properties during quenching. Fig. 2 shows the tensile strengths of the samples with an isothermal holding treatment at different temperatures for different times. The results indicate that the tensile strength decreases rapidly with prolonged isothermal duration at the initial stage, and then the rate of change decreases as the duration increases. At a moderate temperature of 350 C, it takes 30 s to decrease from 480 MPa to 391 MPa, and then it takes 240 s to decrease from 391 MPa to 278 MPa. Furthermore, the rate of change of tensile strength also depends on temperature. Below 350 C, the decreasing rates increase with decreasing temperatures; conversely, the decreasing rate decreases with increasing temperatures from 350 to 480 C. When time is held constant at 300 s and tensile strength is analyzed at 250, 350, and 480 C, it is apparent that the rate of change of the tensile strength is not affected as much at low and high temperatures as it is at the moderate temperature. At 250 C, the initial tensile strength of 480 MPa only decreased to 441 MPa and at 480 C it only decreased to 467 MPa, whereas at the moderate temperature of 350 C, the tensile strength decreased to 278 MPa after 300 s. Fig. 3 shows the yield strengths of the as-aged samples with an isothermal holding treatment at different temperatures for different times. The trend of the yield strengths with temperature and isothermal time is similar to that of the tensile strengths.
3.1. Phase transformation fractions and tensile properties In this work, the time-temperature-transformation fraction diagram was used to show the evolution of the phase transformation during the isothermal holding treatment. The values of the electrical conductivities (g) were transformed to phase transformation volume fractions (f) by using Eq. (1) [15]. In this equation, the value of electrical conductivity [34% (IACS)] of the as-quenched sample
Table 1 Chemical composition of the studied material (wt%). Cu
Mg
Mn
Si
Fe
Ni
Zn
Ti
Al
4.28
0.6
0.81
0.94
0.15
0.003
0.01
0.04
Balance
3.2. TTP diagrams The average tensile strength and yield strength of the test samples without any isothermal holding time were 480 MPa and 430 MPa, respectively. These values were taken as the maximum values, and then the contours of percentage of the tensile properties for each sample as a function of different isothermal treatments were constructed. Fig. 4 shows the TTP contours for the tensile strengths and yield strengths. The results indicate that the shape of the contours is a “C” type. The tensile strengths and yield strengths decrease most rapidly at 350 C. At a given temperature, the breadth of the iso-contours from 100% to 60% of the yield strengths is smaller than that of the tensile strengths. This means that the
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Fig. 1. Electrical conductivities and phase transformation volume fractions of the as-quenched samples with isothermal holding treatment at different temperatures from 250470 C for different times: (a, b) electrical conductivities; (b) phase transformation volume fractions.
Fig. 2. Tensile strengths of the as-aged samples with isothermal holding treatment at different temperatures for different times: (a) 480-350 C; (b) 350-250 C.
yield strengths decrease faster than the tensile strengths with increased isothermal holding times. TTP curves derived from the tensile strengths were used to further investigate the quenching sensitivity of the studied material. The TTP curve can be described by Eq. (2) [16]. By using this
equation, the critical time (tc) to produce a certain percentage of the material at a maximum tensile properties can be calculated, since the parameter tc is related to the amount of quench-induced phase transformation volume fraction.
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Fig. 3. Yield strengths of the as-aged samples with an isothermal holding treatment at different temperatures for different times: (a) 480-350 C; (b) 350-250 C.
Fig. 4. TTP contours for 2A14 aluminum alloys: (a) tensile strengths; (b) yield strengths.
1 tc ¼ k1 k2 exp RT
k3 k24 ðk4 TÞ2
!! þ k5
(2)
where k1 is the natural logarithm of the untransformed fractions; k2 is a constant related to the reciprocal of the number of nucleation sites (s); k3 is a constant related to the energy required to form a nucleus (J$mol1); k4 is a constant related to the solvus temperature; k5 is a constant related to the activation energy for diffusion (J$mol1); R is the gas constant (J$mol1$K1); and T is the absolute temperature (K). According to the tensile strength results, the times to achieve 95% of the maximum tensile strength at different temperatures were used to fit Eq. (2). The coefficients in Eq. (2) were evaluated by an iterative non-linear fitting method. The coefficients k2-k5 are shown in Table 2. By using Eq. (2) and the coefficients in Table 2, the TTP curves of the studied material were derived at 95%, 90%, 85%, and 80% of their maximum levels, as shown in Fig. 5. The nose Fig. 5. Time-temperature-tensile strengths (TTP) curves.
Table 2 Coefficients of Eq. (2). k3/J$mol1
k2/s 1.21 10
12
9.71 10
3
k4/K
k5/J$mol1
1034
108462
temperature is about 350 C. The tensile strength decreases rapidly in the temperature range from 300 to 390 C, while, it changes
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Fig. 6. TEM micrographs of the as-quenched samples without isothermal holding: (a) bright-field TEM; (b) EDX spectra at position A.
Fig. 7. TEM micrographs of the as-quenched samples after isothermal holding for different times at 350 C: (a) 30 s (dark field); (b) EDX spectra at position P1 in (a); (c) 1200 s (bright field); (d) XRD analyses after holding for different times at 350 C.
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Fig. 8. {001}Al TEM micrographs of the as-aged samples after isothermal holding for different times at 350 C: (a, b) 10 s; (b) the SAD pattern of (a); (d) 30 s; (e) 300 s.
slowly in other temperature ranges, especially at temperatures above 450 C. It requires more than 450 s for the tensile strength to decrease to 95% of the maximum value at 480 C, but less than 10 s to decrease to 95% of the maximum value at 350 C. To obtain 95% of the maximum tensile strength for the forged 2A14 aluminum alloy components, the critical average cooling rate was about 7 C s1 at temperatures from 300 to 390 C. 3.3. Microstructural observation In order to clarify the microscopic mechanism of quenching sensitivity, TEM, EDX, and XRD analyses were used to investigate the precipitation behavior of the studied material at 450 C, 350 C, and 250 C with different isothermal holding times. Fig. 6 shows the microstructure of the as-quenched sample without any isothermal holding time. As shown in Fig. 6 (a), there are a few rod-shaped dispersed particles (arrow A). Their lengths are in the range of 100e350 nm and their diameters are in the range of 50e120 nm. According to the EDX results [Fig. 6 (b)], the precipitates are enriched with Cu/Mn. Referring to Ref. [17], it can be concluded that this phase is the T (Al20Cu2Mn3) phase. Except for the T phase, there are no other precipitate phases in the matrix. Fig. 7 shows the microstructure of the as-quenched sample after isothermal holding at 350 C for 30 s and 1200 s. The size and density of the quench-induced precipitate phase increases with increasing isothermal holding time. In the case of the sample with
isothermal holding for 30 s, the T phase in Fig. 7 (a) shows no obvious coarsening, compared to the sample without isothermal treatment [Fig. 6 (a)]. However, a large number of quench-induced rod-shaped precipitates with a length in the range of 70e350 nm (arrow P1) appeared in the matrix. According to the EDX analysis [Fig. 7 (b)], the precipitates are enriched with Cu. Referring to Refs. [12,18], these should be q0 phases. In the case of the sample with isothermal holding for 1200 s, the density and size of the quench-induced rod-shaped phase increases dramatically, compared to the sample with isothermal holding for 30 s, as shown in Fig. 7 (c). The long isothermal holding time caused the sample to lose more solutes, which reduced the volume fraction and density of fine age-induced precipitates and retarded the growth of them in the subsequent aging process. This resulted in poorer mechanical properties, specifically a lower yield strength. The XRD result also shows in Fig. 7 (d) that with the increased isothermal holding times, the density of the quench-induced precipitated phases (Al2Cu) increased, which is consistent with the microstructural observations. Fig. 8 shows the microstructure of as-aged samples with different isothermal holding times at 350 C. In the case of the sample subjected to isothermal holding for 10 s, in addition to the coarse T-phases, another needle-shaped precipitate phase appeared with a length of about 10e100 nm, as shown in Fig. 8 (a). This new phase is distributed homogeneously along the {001}Al planes in the matrix. The corresponding SAD pattern shows that the
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Fig. 9. {001}Al TEM micrographs of the as-aged samples after isothermal holding for different times at 350 C: (a) 0 s; (b) 10 s; (c) 30 s; (d) EDS mapping images for the sample subjected to isothermal holding for 10 s.
streaks are discontinuous [Fig. 8(c)]. Based on prior literature [17,19e21], Fig. 8(a, c) suggests that the age-induced needle-shaped precipitate is a q0 fine precipitate. As shown in the HRTEM image [Fig. 8 (b)], there are some small needle-shaped precipitates along the {001}Al planes in the matrix, comprising a triple copper-layer structure with a length of about 15 nm. According to the literature [22], this should be a q00 fine precipitate, which means that the q0 and q00 precipitated phases existed together in the matrix. With the increase of isothermal times to 30 s, the density of small ageinduced needle-shaped fine q0 phases is obviously reduced and the particles are distributed unevenly in the matrix, as shown in Fig. 8 (d). In addition, the density of coarse rod-shaped q0 phases increases and there is a large precipitate free zone (PFZ) around the coarse phases. When the isothermal time is extended to 300 s, the size of the coarse precipitates (q phases) increases sharply to about 500 nm in length and the fine precipitates almost disappear. Fig. 9 shows precipitated phases along the grain boundaries of the as-aged samples with different isothermal holding times at 350 C. In the case of the sample without any isothermal holding treatment, the rod-shaped precipitates, with a length of about 100 nm, are distributed discontinuously along the grain boundary, and the PFZ zone is smaller than 40 nm. In the case of the sample that was subjected to isothermal holding for 10 s, the precipitated phases along the grain boundary are distributed continuously, and the PFZ zone increases to about 160 nm. For the sample with isothermal holding for 30 s, the precipitates grow to about 200 nm in length, and the PFZ zone increases to more than 240 nm. As shown in Fig. 9 (d), the shapes of the precipitated phases along the grain boundary are rectangular, and they are enriched with Cu. We inferred that the phase is q0 (q) (Al2Cu) in the grain boundaries. Fig. 10 shows the microstructures of the samples subjected to isothermal holding for 30 s at 450 C and at 250 C. In contrast to
the microstructures of the 350 C samples in Figs. 7 (a) and 10 (a) shows that there are no quench-induced rod-shaped precipitated phases, except for the coarse T phase, in the matrix of the asquenched sample isothermally held at 450 C. On the other hand, there are some quench-induced precipitated phases, again except for the coarse T phase, in the matrix of the as-quenched sample isothermally held at 250 C, as shown in Fig. 10 (b). However, the amount is smaller than at 350 C. As shown in Fig. 10 (cef), in the case of the as-aged samples isothermally treated at 450 and 250 C for 30 s, the needle-shaped q0 precipitated phases are smaller in length and distributed uniformly in the matrix, compared to the sample isothermally treated at 350 C for the same time [Figs. 8 (d) and 9 (c)]. In addition, the precipitated phases along the grain boundaries are distributed discontinuously, and the PFZ zone is smaller. The results indicate that the quench-induced phase precipitation rate is much smaller at 450 C and 250 C than at 350 C. Consequently, for the same isothermal holding time, the degree of supersaturation of the as-quenched samples treated at 450 C and 250 C remains high, which increase the density and volume fraction of fine age-induced precipitates after aging treatment, which improves the mechanical performance. This result is in agreement with the results of the mechanical properties tests, as shown in Figs. 2 and 3. 4. Discussion This work shows that for the studied material, increasing the isothermal holding times increases the phase transformation fractions during quenching, and reduces the tensile strengths and yield strengths after aging treatments. These material properties change rapidly during the quenching sensitivity temperature range from 300-390 C, and most rapidly at the nose temperature of 350 C.
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Fig. 10. {001}Al TEM bright field micrographs of the as-quenched and as-aged 2A14 aluminum alloy after isothermally holding for 30 s at 450 C and 250 C: (a) as-quenched sample isothermally held at 450 C; (b) as-quenched sample isothermally held at 250 C; (c, d) as-aged sample isothermally held at 450 C; (e, f) as-aged sample isothermally held at 250 C.
When the quench-induced phase transformation fractions increase, the density of fine age-induced precipitates and the breadth of PFZ zone decrease after the aging treatments. The metals and alloys can be strengthened by the fine secondphase particles, and the strengths increase with the density and size of fine second-phase particles [23,24]. For this studied material, the spaces between the quench-induced precipitates are big, and it is easy for the dislocations pass through it [25]. The fine ageinduced precipitates are the main strengthening particles. The density and volume fraction of the age-induced precipitates decrease with increasing quench-induced phase transformation fraction during quenching. In addition, the breadth of the PFZ zone increases with increasing quench-induced phase transformation fraction. Because the quench-induced phases are coarse and grow during the subsequent aging treatment, which consume the solute atoms in the matrix, especially the atoms around the coarse quench-induced precipitated phases. During isothermal quenching, the phase transformation fraction is determined by the nucleation and decomposition of the supersaturated solid solution. The quench sensitivity is determined by the concentration of alloy elements and nucleation sites [26,27]. The nucleation rate during quenching can be described by Eq. (3) [12,28]:
I ¼ cexpð DG* þ Q =ðRTÞ
(3)
where c is a constant, DG* is the nucleation activation energy, Q is the solute diffusion activation energy, R is the gas constant, and T is the absolute temperature. Eq. (3) shows that the phase-nucleation rate increases as the temperature is lowered. Meanwhile, the diffusion process significantly affects precipitation and growth of the quench-induced phases, and diffusion rate decreases with decreasing temperature. The precipitation processes during quenching are determined by temperature, and the effects are complicated. In general, at high temperatures near the solution temperature, the diffusion rates of solute atoms are high; however, the degree of supersaturation of solute atoms is very low and the nucleation rate is very small. Consequently, the precipitation rate is small. On the other hand, the nucleation rate is high at low temperatures because of the high degree of supersaturation, but the diffusion rate is slow, which results in a low precipitation rate. Meanwhile, at the intermediate temperatures, the precipitation rate is high, because both the nucleation rate and the diffusion rate are high, which explains why the quenching sensitivity temperature range occurs at intermediate temperatures. Even though the evolution of microstructures and tensile properties of the forged 2A14 aluminum alloy during quenching-
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aging treatment has been analyzed here, there are still some insufficiencies. First, this work did not explore why the yield strength decreases faster than tensile strength with increasing isothermal holding time. Second, the influence of heating treatments on the evolution of microstructures and mechanical properties was not investigated. For the studied material, the heating treatment varies for different aims [29,30]. In addition, heating treatments and the amount of plastic deformation affects the quenching sensitivity of aluminum alloys [31,32]. For example, T73 treatment decreases the nose temperature of the TTP curve by 40 C for the 7075 aluminum alloy, compared to T6 treatment [31]. In order to expand the application of this study, the influence of heating treatments on the microstructural evolution and mechanical properties should be investigated in future work. 5. Conclusion In this work, the quenching sensitivity of forged 2A14 aluminum alloy was investigated by the time-temperature-tensile strength/ yield strength diagrams. TEM, EDX, and EDS methods were applied to observe the changes in microstructure. Based on the results and analysis, several conclusions can be drawn: 1) The nose temperature of the TTP contours and curves of the forged 2A14 aluminum alloy is about 350 C, and the quenching sensitivity temperature zone is from about 300 to 390 C. The mechanical properties degrade rapidly in the quenching sensitivity temperature range and most rapidly at 350 C, but the changes became slow above or below this temperature. 2) The heterogeneous precipitation of coarse equilibrium q0 and q phases reduces the concentration of solutes, and they precipitate rapidly at temperatures from 300 to 390 C during isothermal holding treatments, especially at 350 C. The increasing quench-induced phase transformation fractions reduce the density and volume fraction of the fine age-induced needle-shaped q0 and q00 precipitates, and increase the size and the breadth of the PFZ zone, which results in a lower mechanical performance. 3) Yield strengths decrease faster with increasing isothermal holding time than tensile strengths. 4) The critical average cooling rate is about 7 C/s in the quenching sensitivity temperature range to obtain 95% of the maximum tensile strength. 5) In order to balance the residual stress and mechanical performances of the forged 2A14 aluminum alloy components, it is suggested that the cooling rate should be enhanced in the temperature range from 300 to 390 C and suppressed in the other temperature ranges. Acknowledgments This work was financially supported by State Key Laboratory of High Performance Complex Manufacturing (zzyjkt2014-02) and the fund of Jiangsu Province for the transformation of scientific and technological achievements (BA2015075). References [1] Y. Xiao, S. Xie, J. Liu, T. Wang, Practical Handbook of Aluminum Technology, Metallurgical industry Press, Beijing, 2005. [2] B. Milkereit, M.J. Starink, Quench sensitivity of Al-Mg-Si alloys: a model for linear cooling and strengthening, Mater. Des. 76 (2015) 117e129. [3] B.C. Shang, Z.M. Yin, G. Wang, B. Liu, Z.Q. Huang, Investigation of quench sensitivity and transformation kinetics during isothermal treatment in 6082
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