Characteristic constitution model and microstructure of an Al-3.5Cu-1.5Li alloy subjected to thermal deformation

Materials Characterization 145 (2018) 53–64

Contents lists available at ScienceDirect

Materials Characterization journal homepage: www.elsevier.com/locate/matchar

Review

Characteristic constitution model and microstructure of an Al-3.5Cu-1.5Li alloy subjected to thermal deformation

T

⁎

Weichen Yua, Hongying Lia,c,d, , Rong Dub, Wen Youb, Mingchun Zhaoa, Zheng-an Wangb a

School of Materials Science and Engineering, Central South University, Changsha 410083, China Southwest Aluminum (Group) Co., Ltd., Chongqing 401326, China c State Key Laboratory on Lightweight High-strength Structural Material, Changsha 410083, China d Nonferrous Metal Oriented Advanced Structural Materials and Manufacturing Cooperative Innovation Center, Central South University, Changsha 410083, China b

A R T I C LE I N FO

A B S T R A C T

Keywords: Al-Li alloy Thermal deformation Constitution model Dynamic precipitation Continuous dynamic recrystallization

The constitution model and microstructure of 2A97 Al-Li alloy during thermal deformation were investigated. The thermal deformation behavior was dominated by the interaction between work hardening and dynamic softening. The constitution model was constructed to predict the rheology behavior, in which Zener-Holloman equation with a fifth-order polynomial was set to describe the effect of various strains on the material parameters. Processing maps displayed to exist two unstable domains in microstructure (one domain with the temperature range from 300 °C to 370 °C and the strain rate range from 100.5 s−1 to 10s−1, and the other with the temperature range from 380 °C to 500 °C and the strain rate range from 0 s−1 to 10 s−1). Microstructural characteristics of dynamic recovery (DRV) and dynamic recrystallization (DRX) in the stable domain and local rheology in the unstable domain were observed, respectively. The dynamic precipitation (DPN) of T1 phases and the interaction between dislocations and Al20Cu2Mn3 and Al3Zr phases were concluded by high resolution transmission electron microscopy (HRTEM), which was the first experimental evidence for the DPN in 2A97 AlLi alloy. All these confirmed the occurrence of the continuous dynamic recrystallization during thermal deformation.

1. Introduction The conventional aluminum alloys are being replaced gradually by aluminum‑lithium (Al-Li) alloys in the aerospace industry because Al-Li alloys have low density, high specific stiffness, high specific strength, good corrosion resistance, excellent low-temperature performance and superplastic forming performance. So far, the rapid developments of aircraft rockets, missiles, military aircrafts and modern airliners have promoted the industrialization of the third generation Al-Li alloys such as 2198-T8 [1] and 2060-T8E30 [2]. By increasing the amount of copper as well as other alloy elements and reducing the amount of lithium, the rigidity was increased by 20% and the elastic modulus was increased by 15% without loss of tensile strength and corrosion resistance [3,4]. As the key structural material for the cabin sheet and aircraft wing skin, Al-Li alloys are required to have good workability, whose preparation need a thermal plastic deformation after the homogenization of casting ingots. The ranges of temperatures and strain rates are narrow during thermal deformation, which will lead to the high

⁎

cracking risk during hot rolling because of strong local deformation tendency. Due to the poor hot workability, 2A97 Al-Li alloy requires suitable hot working processes to reduce the heterogeneity of microstructures and properties for the thick plates and sheets. Many efforts have been focused on the rheological behaviors of aluminum alloys [5–7]. Several empirical models based on phenomena and physics such as Jackson-Cook (JC) model and Back PropagationArtificial Neural Network (BP-ANN) model have been widely used to describe the thermal deformation behaviors [8]. Recently, JC model with the temperature raising and friction correction was successfully used under the given strain [9]. However, the strains are always changed during the whole thermal deformation processing and these models are not appropriate if the effect of various strains is considered. On the other hand, the addition of Mg could reduce the stacking fault energy of aluminum alloys and result in the occurrence of the recrystallization during the thermal deformation of Al-Mg [10], Al-ZnMg-Cu [11] and Al-Cu-Mg-Li [12] alloys. The nucleation and the growth of the dynamic recrystallization (DRX) grains were detected in and Al-Mg-Li-Zr [13] alloys. Moreover, the dynamic precipitation

Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha 410083, China. E-mail address: [email protected] (H. Li).

https://doi.org/10.1016/j.matchar.2018.08.029 Received 17 April 2018; Received in revised form 11 July 2018; Accepted 17 August 2018 Available online 18 August 2018 1044-5803/ © 2018 Published by Elsevier Inc.

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 1. The true stress-strain curves of tested specimens at (a) deformation temperatures on 400 °C with various strain rates and (b) strain rate on 1 s−1 with various deformation temperature.

Fig. 2. Relationship and linear fit between (a) lnσ–lnε ̇ and (b) σ–lnε .̇

Fig. 3. Relationship and linear fit between (a) lnε –̇ ln[sinh(ασ)] and (b) 1/T–ln[sinh(ασ)].

2. Experimental Procedures

(DPN) occurred during hot working in AA2195 [14] and 2050 alloys [15]. However, the mechanism and experimental evidence of DRX in 2A97 Al-Li alloy is not known nor is the occurrence of DPN in 2A97 AlLi alloy because high alloying addition make them much more complex. In the present work, the thermal deformation behaviors of 2A97 AlLi alloy were investigated systematically by true stress-strain curves. The characteristic constitution model was established to describe the effect of various strains on rheology behavior. Processing maps were drawn to predict the plastic deformation mechanism and the unstable domains under various deformation conditions. Microstructures during thermal deformation were analyzed to illuminate the mechanism of dynamic recrystallization.

A commercial 2A97 Al-Li alloy with the chemical composition of Al3.5Cu-1.5Li-0.5Mg-0.5Zn-0.3Mn-0.1Zr-0.08Ti was detected for the present researches. The ingots were smelted in a vacuum melting furnace by high purity aluminum, lithium as well as other master alloys under dynamic Argon atmosphere, then cast into rectangular ingots (160 mm × 240 mm × 35 mm). The ingot was homogenized at 410 °C for 8 h followed by 500 °C for 24 h in salt bath furnace. Cylindrical specimens with the size of 10 mm diameter and 15 mm height were cut from the central part of the ingots subsequent to homogenization. The isothermal compression test was performed using a Gleeble-3500 54

Materials Characterization 145 (2018) 53–64

W. Yu et al.

is mainly due to the increase of dislocation movement resistance [17]. Therefore, the effect of work hardening is much greater than that of dynamic softening in the early stage of deformation. In the middle stage of the deformation, as the strain increases gradually, the rate of increase in flow stress gradually slows down so the work hardening and dynamic softening reach equilibrium and flow stress achieves a peak stress in Fig. 1(a). In the later stage of deformation, the degree of dynamic softening continues to increase and exceeds the degree of work hardening, leading to the decrease of flow stress. And then work hardening and dynamic softening tend to be dynamic equilibrium and the flow stress is stabilized gradually. The results of the true stress-strain curves under other deformation conditions are similar with those of Fig. 1(a) and (b). It is worth noting that at low strain rates (≤0.1 s−1), the dynamic softening and work hardening are dynamically balanced as the strain continues to increase, and the flow stress remains essentially constant. However, at high strain rates (1–10 s−1), the softening mechanism is mainly dominated by DRX and the flow stress decreases continuously. Fig. 4. Relationship and linear fit between ln[sinh(ασ)] and lnZ.

3.2. Establishment of Constitutive Equation Table 1 Values of material constants β, n, α, Q and lnA about strain.

β n α Q lnA

0.1

0.2

0.3

0.4

0.5

0.6

0.07841 6.1504 0.01248 167.0901 29.4026

0.07594 6.1743 0.01230 166.4618 28.6977

0.07411 6.1206 0.01210 165.8800 27.5619

0.07352 6.0552 0.01214 163.5275 26.6533

0.07221 5.7629 0.01253 162.2561 26.2059

0.07443 0.5779 0.1288 160.0661 26.1124

Based on the true stress-strain curves, the relationship between flow stress and deformation parameters such as temperature and strain rate can be determined by constructing the constitutive equation [18]. The Arrhenius model has been successfully used to predict the flow stress [19], which can be described as Eq. (1).

Z = ε ̇ exp(Q/ RT )

(1)

Under different stress levels, the relationship between different deformation parameters can be expressed by Eqs. (2) and (3) [18,20].

thermal mechanical simulator at the temperature range from 300 °C to 500 °C with a 50 °C interval and four different strain rates from 0.01 s−1 to 10 s−1. The specimens were heated to the given temperature at a rate of 10 °C/s and soaked for 5 min to uniform temperature distribution. After that, they were deformed to 40% of the original height and then quenched in water to room temperature immediately. Microstructures such as grain boundary angle orientation, deformation microstructure distributions and recrystallization of the samples were analyzed by using Helios Nanolab-600i dual beam electron microscopy (FIB/SEM) device equipped with an HKL Channel 5 EBSD system. The tiny sub-structures and precipitates were observed by Titan G2 60–300 transmission electron microscope (TEM) with the acceleration voltage of 200 kV. The specimens were mechanically thinned to 0.08–0.1 mm and polished by twin jet electropolishing system, where the mixture of 25% HNO3 + 75% CH3OH (volume fraction) was selected as electrolyte. The voltage and the electric current were set at 15 V and 10–20 mA respectively. The temperature was controlled below −30 °C during the operation.

When ασ < 0.8, ε ̇ = A1 σ n1 exp(−Q/RT )

(2)

When ασ > 1.2, ε ̇ = A2 exp(βσ ) exp(−Q/ RT )

(3)

For all stress levels, the relationship is as followed:

ε ̇ = A [sinh(ασ )]n exp(−Q/ RT )

(4)

where σ is the flow stress corresponding to a given strain, ε ̇ is the strain rate, R is the gas constant, α is experimentally determined constant of stress, n is the stress exponent, Q is the deformation activation energy, T is the absolute temperature, Z is the Zener-Holomon parameter, A1, A2, A, β, and n1 are all temperature-independent material constants, and

α = β / n1

(5)

Combining Eqs. (1) and (4):

Z = ε ̇ exp(Q/RT ) = A [sinh(ασ )]n

(6)

and

sinh−1 (ασ ) = ln{(ασ ) + ((ασ )2 + 1)1/2}

(7)

So, the flow stress can be expressed as Eq. (8) with Z constant.

3. Results and Discussion

σ = (1/ α ) ln{(Z / A)1/ n + ((Z / A)2/ n + 1)1/2} 3.1. The True Stress-Strain Curves

(8)

The logarithms to Eqs. (2) and (3) are as followed:

ln ε ̇ = ln A1 + n1 ln σ − Q/ RT

The true stress-strain curve of samples compressed at deformation temperatures of 400 °C with various strain rates was shown in Fig. 1(a). The curve of samples compressed at strain rate of 1 s−1 with various deformation temperatures was in Fig. 1(b). It can be seen from Fig. 1(a) and (b) that the flow stress is sensitive to deformation temperatures and strain rates, which decreases with the increase of the deformation temperature and increases with the increase of the strain rate. The thermal deformation behavior of the alloy is dominated by the interaction between work hardening and dynamic softening (dynamic recovery (DRV) and dynamic recrystallization (DRX)) [16]. The thermal deformation process can be divided into three stages according to Fig. 1. In the early stage of deformation, the rapid increase in flow stress

ln ε ̇ = ln A2 + βσ − Q/ RT

(9) (10)

From Eqs. (9) and (10), the effect of strain on flow stress is not considered. However, it was illustrated in present work that the deformation activation energy and the material constants of the alloy may be changed due to the change of microstructural evolution mechanism of the material deformation with the increase of strain [21]. Therefore, the material constants of the tested specimens were calculated in the case of strain 0.1. Fig. 2 shows the relationship and linear fit between lnσ–lnε ̇ and σ–lnε .̇ The average slope of the fit lines shown in Fig. 2(a) is n1, which is 55

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 5. Relationship and linear fit between true strain and (a) β, (b) α, (c) n, (d) Q and (e) lnA. Table 2 Polynomial function coefficient of material constants β, α, n, Q and lnA. β

α

B0 B1 B2 B3 B4 B5

0.06959 0.23028 −2.04571 7.37875 −12.02917 7.30833

C0 C1 C2 C3 C4 C5

n 9.8714 −86.73617 730.3175 −2825.2875 5094.75 −3474.58333

Q −0.10334 2.64524 −21.71825 81.96792 −144.525 96.30833

D0 D1 D2 D3 D4 D5

calculated as 7.9467. The average slope of the fit lines shown in Fig. 2(b) is β, which is calculated as 0.07841. α is equal to β divided by n1 according to Eq. (5), so α is calculated to be 0.009867. The logarithm to Eq. (4) is as followed:

ln ε ̇ = ln A + n ln[sinh(ασ )] − Q/ RT

lnA

E0 E1 E2 E3 E4 E5

183.7716 −354.95142 2650.03458 −8991.97083 13,846.04167 −7933.75

∂ ln ε ̇ ⎫ ⎧ ∂ ln(sinh(ασ )) ⎫ Q = R⎧ ⎨ ⎬ ∂ (1/ T ) ⎩ ∂ ln(sinh(ασ )) ⎬ ⎭T ⎨ ⎩ ⎭ε ̇ Set

{

F0 F1 F2 F3 F4 F5

28.5116 22.20568 −164.72792 345.04167 −280.20833 69

(12)

}

to be M, which is the representative of the average

}

to be N, which is the representative of the average

∂ ln ε ̇ ∂ ln(sinh(ασ )) T

slope on the fit lines of the relationship between ln[sinh(ασ)] and lnε ,̇

(11)

and set The differential equation of Eq. (11) is as followed:

{

∂ ln(sinh(ασ )) ∂ (1 / T ) ε̇

slope on the fit line of the relationship between 1/T and ln[sinh(ασ)]. 56

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 6. Processing maps of the tested alloy with different strain: (a) 0.1, (b) 0.2, (c) 0.3, (d) 0.4, (e) 0.5 and (f) 0.6.

calculated as 167.0901 kJ/mol. According to the previous research, the activation energy for lattice self-diffusion in aluminum was calculated as 1.45 ± 0.05 eV (≈142.31 kJ/mol) [22]. Since Q value of 167.0901 kJ/mol is higher than the activation energy for lattice selfdiffusion in aluminum, which indicated the deformation mechanism related to dislocation climb and cross-slip during hot working. Combining Eqs. (1) and (4) and taking the logarithm, then:

Table 3 The range of temperature, strain rate and power dissipation efficiency of domains I–IV. Domain I II III IV

Temperature/°C

Strain rate/s−1

η/%

300–370 470–500 300–365 380–500

10 –10 10−2–10−0.4 10–1.75–10−0.5 0–10

16–30 24–32 16–26 14–22

0.5

ln Z = ln A + n ln[sinh(ασ )]

(13)

Fig. 4 shows the relationship and linear fit between ln[sinh(ασ)] and lnZ, in which the intercept and slope of the fit lines are 29.4026 and 6.1504 respectively. By replacing the value of n1 with the value of the slope in Fig. 4, the material constants can be calculated more accurately and reliably. So, lnA is equal to 29.4026 and A can be calculated as 5.88 × 1012 s−1.

So, Q is equal to R multiplied by M multiplied by N. Fig. 3 shows the relationship and linear fit between ln[sinh (ασ)]–lnε ȧ nd 1/T–ln[sinh(ασ)]. The average slope of the fit lines shown in Fig. 3(a) is calculated as 3.6969, which shown in Fig. 3(b) is calculated as 5.4359. So, the deformation activation energy Q can be 57

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 7. Inverse pole figures (left) and deformation microstructure distributions (right) of the tested alloy at strain rate of 10 s−1 with different temperatures of (a, b) 300 °C, (c, d) 350 °C, (e, f) 400 °C and (g, h) 500 °C.

According to the steps above, the corresponding flow stress values under different strains (0.1–0.9) are brought into the solution procedure and the material constants (β, n, α, Q and lnA) of the alloy under various strain are obtained, which is shown in Table. 1. It can be seen from Table 1 that the strain has a significant effect on

On the basis of the above analysis, a constitutive equation that describes the flow stress of the tested alloy as a function of ε ̇ and T at ε = 0.1 can be written as:

ε ̇ = 5.88 × 1012 [sinh(0.0099σ )]7.95 exp( −167090/ RT )

(14) 58

Materials Characterization 145 (2018) 53–64

W. Yu et al.

[27]. However the highest value of η does not mean the best workability because unstable deformation (such as the formation of a crack) also bring about a large power dissipation efficiency [28]. According to the principle of irreversible thermodynamics in metal materials during the severe plastic deformation process, the criterion for the flow instabilities can be descripted as follows [25]:

ξ (ε )̇ =

the material constants. The constitutive equation constructed by materials constants expressed by the polynomials of the strain can reflect the relationship between the stress and deformation parameters, which is more consistent with the actual deformation conditions and the results are more accurate and reliable [23]. The relationship and linear fit between the strain and material constants in Fig. 5 shows that β, n, α, Q and lnA can be expressed by the fifth-order polynomial of the strain respectively, in which the correlation coefficient is up to 99.82%. Combining Eq. (8) and the fifth-order polynomial between the strain and material constants, the constitutive equation is expressed as Eq. (15), where the coefficients are indicated in Table 2. After all, the characteristic constitution model based on Arrhenius model was similar to those of 2099 alloy subjected the thermal deformation, where the activation energy Q was 203.249 kJ/mol and the stress exponent n was 5.0547 [24]. These values indicated grain rotation associated with dislocation climb and cross-slip governed plastic flow.

(15)

3.3. Processing Map The processing maps of the tested alloy were built based on the dynamic material modeling (DMM), which is constructed with modeling of physical systems and irreversible thermodynamic principles [25]. The processing maps are superimposed on the power dissipation efficiency maps and instable domain maps, which can reflect the relationship between the microstructure evolution mechanism and the thermal deformation parameters of the alloy, and indicate the stable domains and unstable domains throughout the thermal deformation [26]. The proportion of the total power dissipated by the microstructure evolution of the deformed alloys is called the efficiency of power dissipation, η, which is a dimensionless parameter given as:

η=

2m m+1

(17)

where ξ (ε )̇ is a dimensionless instability parameter varied with temperature and strain rate. In the processing maps, the region where ξ (ε )̇ is negative represents flow instability. The processing maps at different true strains (0.1–0.6) were plotted in Fig. 6, where the values in the contours of the maps represent the percentage of power dissipation efficiency η with the given strain, the shaded parts represent the unstable domains and the rest parts are recognized as the stable domains. As the strain being 0.1 in Fig. 6(a), the value of η is higher than 30% in two main domains. One is a low temperature domain in the temperature range of 320–360 °C and strain rate range of 100.7–10s−1, and another is a high temperature domain in the temperature range of 480–500 °C, and strain rate range of 10−2–10−1.5 s−1. There exists an unstable domain in the temperature range of 300–400 °C and the strain rate range of 10–1.835–10−0.58 s−1, where the peak value of η is 25%. When the strain is 0.2, there is a decrease of η in the temperature range of 480–500 °C and the strain rate range of 10−2–10−1.5 s−1. The temperature range of the unstable domain becomes 300–390 °C in Fig. 6(b). When the strain is 0.3, the area of the low temperature unstable domain continues to decrease, leading to the temperature range of 300–380 °C. In the temperature range of 400–500 °C and strain rate range of 100.6–10s−1, there appears a second unstable domain zone, where the power dissipation efficiency η is from 20% to 25%. As the strain being 0.4, the spacing of the contour of η becomes narrower and the peaks of the value of η becomes more obvious, which can be up to 32% in Fig. 6(d). The area of the low temperature unstable domain continues to decrease, where the temperature range is reduced to 300–370 °C. However, the area of the high temperature unstable domain increases, where the strain rate range becomes 100.45–10s−1. When the strain becomes 0.5 and 0.6 in Fig. 6(e) and (f), the value of η increases slightly with the increase of strain. The peak of η is mainly concentrated at two domains. One is in the temperature range of 320–360 °C and the strain rate range of 100.85–10s−1, and the other is in the temperature range of 480–500 °C and the strain rate range of 10–1.75–10−0.5 s−1. The area of low temperature unstable domain is decreasing as the temperature range only being 300–365 °C and the strain rate range being 10–1.75–10−0.5 s−1. However, the area of high temperature unstable domain is increasing as the temperature range becoming 380–500 °C and the strain rate range becoming 0–10s−1. Therefore, there are four typical domains in Fig. 6(d), where domains I and III are stable, and domains II and IV are unstable. The range of temperature, strain rate and power dissipation efficiency of domains I–IV are displayed in Table 3. Thermal deformation in the unstable domain may bring about adiabatic shear band, local flow and dynamic strain aging, which are not advantage for the ultimate performance of the alloy [29]. It is necessary to avoid thermal deformation in the conditions corresponding to the unstable domain. In previous work, the peak power dissipation efficiency of DRV is about 23% and that of DRX is about 42% for the tested alloy [30]. It can be seen from the contours in Fig. 6(f) that the peak value of η is 32%, which indicated that DRV and DRX occurred simultaneously under experimental deformation conditions. The dynamic softening of the alloy is favorable for processing of the thermal deformation.

Fig. 8. Deformation microstructure fraction of the tested alloy at strain rate of 10 s−1 with various.

1/ n 2/ n 1/2 ⎧ σ = (1/ α ) ln{(Z / A) + ((Z / A) + 1) } ⎪ ε ̇ = A [sinh(ασ )]n exp(−Q/ RT ) ⎪ ⎪ Z = ε ̇ exp(Q/ RT ) ⎪ β = B0 + B1 ε + B2 ε 2 + B3 ε 3 + B4 ε 4 + B5 ε 5 ⎨ n = C0 + C1 ε + C2 ε 2 + C3 ε 3 + C4 ε 4 + C5 ε 5 ⎪ 2 3 4 5 ⎪ α = D0 + D1 ε + D2 ε + D3 ε + D4 ε + D5 ε ⎪Q = E0 + E1 ε + E2 ε 2 + E3 ε 3 + E4 ε 4 + E5 ε 5 ⎪ 2 3 4 5 ⎩ ln A = F0 + F1 ε + F2 ε + F3 ε + F4 ε + F5 ε

∂ ln(m / m + 1) +m<0 ∂ ln ε ̇

3.4. Microstructures

(16)

Generally, a large value of η means a good workability of the alloys

The 59

Inverse

pole

figures

and

deformation

microstructure

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 9. Inverse pole figures (left) and deformation microstructure distribution maps (right) of the tested alloy at (a) 300 °C with the strain rate of 0.1 s−1, (b) 300 °C with the strain rate of 1 s−1 and (c) 500 °C with the strain rate of 0.1 s−1.

distributions of the tested alloy at strain rate of 10 s−1 with various deformation temperatures are illustrated in Fig. 7. Those deformation conditions are taken from domains I and IV respectively in Table 3. The deformation temperature of 300 °C and 350 °C represent stable domain I and that of 400 °C and 500 °C represent unstable domain IV. It is obvious that there are different microstructure characterizations in different processing domains, including fine sub-grains, recrystallized grains and local deformation banding. Compressing the alloys at 300 °C, grains are elongated, while most of the grain boundary was sharp. A large amount of plastic work from high-speed deformation spread too slowly in a short time to transform into the deformation heat, which results in a rapid increase in the local temperature of the alloy and leads to fine sub-grains appeared near the grain boundaries in Fig. 7(a). The power dissipation efficiency is 16%, indicating the occurrence of DRV. With the deformation temperature increasing to 350 °C, the power dissipation efficiency is increased to 30%. The grain size of the subgrains and the dynamic recrystallization grains near grain boundaries increases, accompanied by the occurrence of the recrystallization grains in the deformation grains. As the temperature continues to increase, the deformation conditions transformed into the unstable thermal

deformation. From Fig. 7(e) and (d), there appears the local deformation banding in the unstable domain, that is to say, a narrow strip of tissue occurs inside the grains of the alloy, which is formed by a large number of fine sub-grains. Local rheology is mainly present in some grains, in which the crystal orientation is different from those of the surrounding grains. With the increase of temperature, the local deformation banding is more obvious. Contrasting the microstructure characteristics of different processing domains, the deformation structure of the alloy is more uniform and the power dissipation efficiency range is from 16% to 30% in the stable domain, indicating the characteristic of DRV and DRX obviously [31]. And with the increase of temperature, the mechanism of microstructure evolution is gradually dominated by DRX. According to the deformation microstructure distributions in Fig. 7(b), (d), (f) and (h), the microstructure is mainly conducted by deformed, recrystallized and substructured (between deformed and recrystallized) organizations. When the strain rate is 10 s−1 and the deformation temperature is 300 °C, the alloy is almost entirely deformed with a small amount of recrystallized distribution in the subgrains. As the temperature increases, recrystallization grain grows and 60

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Schmitt factor than its surrounding grain orientation [32]. The slip systems in those grains is more easily activated during the deformation, leading to the flow localization [33]. The alloy in the stable domain with the increase of the strain rate to 1 s−1 shown was shown in Fig. 9(c), where the power dissipation efficiency is low. The vacancy migration and dislocation climb result in the polygonization and cell structure with a large number of fine sub-grains, which is similar to the microstructure in Fig. 7(a). The orientation and pole distributions are investigated with the temperature of 500 °C and strain rate of 0.1 s−1 in Fig. 9(e). The result shows that with the given deformation condition, the large sized sub-grains and recrystallized grains can be observed in and around the deformation grains. The deformation microstructure distributions were investigated in Fig. 9(b), (d) and (f), which shows that the alloy is basically deformed with the low temperature of 300 °C. The sub-structure decreases with the increase of strain rate. The substructure occurs first in the grain boundary at the temperature of 500 °C and the strain rate of 0.1 s−1. The statistical histogram of the deformation microstructure fraction under the corresponding deformation conditions was illustrated in Fig. 10. The degree of DRX during hot deformation can be represented by the proportion of recrystallized and substructured grains, which is 31% at 300 °C/0.1 s−1. However, the value decreases to 12% at 300 °C/ 10 s−1, indicating that lower strain rate is beneficial to DRX. The substructures account for a great proportion in the microstructure compressed at 500 °C/0.1 s−1 due to the lack of driving force of fully DRX as the strain being 0.6. Therefore, the microstructure evolution in the stable domain is in charge of the interaction of DRV and DRX. Both higher deformation temperature and lower strain rate are helpful for DRX. On the other hand, the microstructure evolution in the unstable domain is mainly dominated by local rheology as the dynamic softening occurring first in the unstable structure. In the present work, three mechanisms for the dynamic recrystallization evolution of aluminum alloys with high stacking fault energy were investigated: discontinuous dynamic recrystallization, that is, traditionally, including the nucleation and growth process of dynamic recrystallization, continuous dynamic recrystallization, including transformation of low angle grain boundary (LAGB) to high angle grain boundary (HAGB) without obvious nucleation and growth of recrystallized grains and geometric dynamic recrystallization, that is, the process of splitting out new recrystallized grains from the deformed grains [34–36]. At the initial stage of deformation, the strain produced by plastic deformation induced the dislocation multiplication and dislocation tangle, which are then moved back to the original grain boundary by DRV, or transform to dislocation wall and sub-grain boundaries throughout polygonization, resulting in an increase of the proportion of LAGBs. The grain boundary orientation fraction of the alloy under various deformation conditions was investigated in Fig. 11. Generally, the grain boundaries with adjacent grains below 10° are represented for LAGB, while the grain boundaries with adjacent grains over 15° are HAGB [37,38]. The LAGB fraction reached to 87.3% with the deformation temperature of 500 °C and the strain rate of 0.1 s−1, however the fraction at 300 °C was just 46.1%, indicating that the dislocations tend to move under higher temperature and accelerate the formation of cell structures, which increasing the LAGB fraction. On the other hand, the LAGB fraction was investigated as 65.8% at 500 °C with the strain rate of 10 s−1. The result shows that with the decrease of the strain rate, the degree of DRX increases, which is consistent with the microstructure evolution in Figs. 7 and 9. In the middle stage of deformation, a large number of dislocations move to the subgrain boundary constantly due to the dislocations continuously induced by strain and DRV, resulting in a rotation of LAGBs to HAGBs. The grain and grain boundary structure of the alloys compressed at 300 °C/10s−1, 300 °C/0.1 s−1 and 500 °C/0.1 s−1 are observed by TEM in Fig. 12. The microstructure at 300 °C/10s−1 consists of dislocations and sub-grain in Fig. 12(a), indicating the

Fig. 10. Deformation microstructure fraction of the tested alloy at various deformation conductions.

Fig. 11. Frequency of misorientation of the test alloy in various deformation conditions.

some sub-structures appear. In the unstable domain, deformation organization is mainly concentrated near the grain boundary with most of the sub-structures inside the grain, which indicates a more obvious trend with the increase of deformation temperature in Fig. 7(h). The statistical histogram of the deformation microstructure fraction under the corresponding deformation conditions was shown in Fig. 8, which indicates that when the strain rate is up to 10 s−1, the grains are almost totally deformed. As the deformation temperature is low, the mechanism of microstructure evolution was mainly dominated by DRV, leading to deformed microstructure to 90%. DRX gradually occurs with the increase of deformation temperature, which results in the decrease of deformed organizations and the increase of recrystallization grains. However, due to the high strain rate, the alloy does not have sufficient time for DRV and DRX. Therefore, as the temperature increases, there appears a large number of sub-structures between the deformed and recrystallized structures, and with deformation temperature of 500 °C, the fraction of sub-structures is close to 40%. The Inverse pole figures and deformation microstructure distributions of the tested alloy with the deformation temperatures in the stable domain II and unstable domain III are illustrated in Fig. 9. The alloy in the unstable domain with the deformation temperature of 300 °C and strain rate of 0.1 s−1 was shown in Fig. 9(a), and the local deformation banding can be observed, where the grain orientation has a larger 61

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 12. TEM images of test alloys at (a) 300 °C with the strain rate of 10 s−1, (b) 300 °C with the strain rate of 0.1 s−1 and (c) 500 °C with the strain rate of 0.1 s−1.

Fig. 13(b) taken along the [110] zone axis was investigated in Fig. 13(d). The lattice parameter of this phase measured directly from the inverse FFT image reveals: an ≈0.409 nm, which is slightly different from the lattice parameter of β' phase (Al3Zr) and the FFT is seen to match well with the patterns taken from the β' phase. The HRTEM image of the polygonal phase in Fig. 13(c) was shown in Fig. 13(e). The chemical composition investigated by EDX shows Al of 86.98%, Mn of 6.28% and Cu of 6.74%, indicating the Al20Cu2Mn3 phase according to the report [39]. This is a phase that always exists throughout the thermal deformation and the interface between this phase and Al matrix is along {110}, which is well matched with Fig. 13(e). Fig. 13(f) shows the TEM image of the test alloy deformed at 500 °C/0.1 s−1, where β' phase and Al20Cu2Mn3 phase were still present, but no precipitation of T1 phase was observed. This is mainly caused by lack of driving force used for forming T1 phase, which was used for forming high angle recrystallized grains throughout DRX with the higher temperature and lower strain rate. At the same time, the low orientation sub-grain boundary was rotated through 120° and grain boundaries cleared gradually. In the last stage of deformation, the sub-grain boundary is in contact with the original grain boundary, resulting in the layered structures formed by deformation grains. The original boundary in the triple junction grain boundary tends to move toward the inside of the deformed grains under the surface tension of the sub-grain boundary and interfacial energy. As the strain increases continuously, grain spheroidization and growth gradually begin through the Y-type mechanism

interaction of DRV and DRX throughout the deformation. With the decrease of strain rate, the dislocation tangle as the original of cell structures can be observed at 300 °C/0.1 s−1 as well as the sub-grain in Fig. 12(b), revealing that lower strain rate can improve DRX. With the increase of deformation temperature to 500 °C, the sub-grain as well as triple junction grain boundaries can be seen obviously in Fig. 12(c), where the recrystallized grain boundaries are straight and clear, showing typical characteristics of DRX. The dynamic restoration is the interaction of DRV and DRX throughout the deformation, in which DRX plays a larger role at higher temperature with lower strain rate. The dynamic evolution of precipitates during deformation was investigated by TEM and HRTEM in Fig. 13. At deformation temperature of 300 °C and strain rate of 0.1 s−1, coarse T1 phases can be observed obviously in Fig. 13(a). However, dislocations are formed inside the grains due to the low temperature and coarse T1 phase is refined and precipitated along the grain boundary and sub-grain boundary. Besides, a strong interaction between the coarse T1 phase and dislocations results in a break of T1 phase. Moreover, the increase of distortion energy caused by deformation together with the increase of surface energy caused by break of T1 phases finally leads to the increase of free energy on T1 phases, which is the reason of T1 phase re-dissolution. There is a depression zone on T1 phase in Fig. 13(a), which might be a re-dissolution contrast banding produced by broken T1 phases or T1 phases cut off by dislocations. On the other hand, the tangles between two precipitates and dislocations were observed in Fig. 13(b) and (c). The HRTEM image and corresponding FFT of the spherical phase in 62

Materials Characterization 145 (2018) 53–64

W. Yu et al.

Fig. 13. TEM images of test alloys: (a)–(c) BF image at 300 °C with the strain rate of 0.1 s−1; (d) HRTEM image and corresponding FFT (inset) of the phase in (b); (e) HRTEM image and corresponding FFT (inset) of the phase in (c); (f) BF image at 500 °C with the strain rate of 0.1 s−1.

polynomial with the correlation coefficient up to 99.82%. 2. The peak value of power dissipation efficiency was 32%, indicating that DRV and DRX occurred simultaneously. The dynamic softening of the alloy was favorable for processing of the thermal deformation. 3. The dynamic precipitation of T1 phase was inferred from the redissolution contrast banding at deformation temperature of 300 °C. It is not beneficial to the dynamic precipitation of T1 phase at higher temperatures and lower strain rates due to lack of driving force, which was used for forming high angle recrystallized grains

after the contact between the upper and lower ends of the sub-grain boundary. This results in the HAGBs and finally new recrystallized grains. 4. Conclusions 1. The constitution model was constructed to predict the rheology behavior, in which Zener–Holloman equation was set to describe the effect of various strains on the material parameters as a fifth-order 63

Materials Characterization 145 (2018) 53–64

W. Yu et al.

throughout DRX. 4. The interaction between dislocations and Al20Cu2Mn3 and Al3Zr phases as well as misorientation of grain boundaries confirmed the occurrence of the continuous dynamic recrystallization during thermal deformation.

isothermal compressed Al–Cu–Li alloy, Mater. Sci. Eng. A 586 (2013) 115–122. [18] B. Li, Q. Pan, Z. Yin, Characterization of hot deformation behavior of as-homogenized Al–Cu–Li–Sc–Zr alloy using processing maps, Mater. Sci. Eng. A 614 (2014) 199–206. [19] G. Chen, L. Chen, G. Zhao, C. Zhang, W. Cui, Microstructure analysis of an Al-Zn-Mg alloy during porthole die extrusion based on modeling of constitutive equation and dynamic recrystallization, J. Alloys Compd. 710 (2017) 80–91. [20] C. Bruni, A. Forcellese, F. Gabrielli, Hot workability and models for flow stress of NIMONIC 115 Ni-base superalloy, J. Mater. Process. Technol. 125–126 (2002) 242–247. [21] P. Li, H. Li, L. Huang, X. Liang, Z. Zhu, Characterization of hot deformation behavior of AA2014 forging aluminum alloy using processing map, Trans. Nonferrous Metals Soc. China 27 (2017) 1677–1688. [22] S.L. Robinson, O.D. Sherby, Activation energy for lattice self-diffusion in aluminum, Phys. Status Solidi 1 (1970) 119–122. [23] J. Guo, S. Zhao, R. Murakami, R. Ding, S. Fan, Modeling the hot deformation behavior of Al alloy 3003, J. Alloys Compd. 566 (2013) 62–67. [24] F. Zhang, J. Shen, X. Yan, J. Sun, X. Sun, Y. Yang, Y. Liu, High-temperature flow behavior modeling of 2099 alloy considering strain effects, Trans. Nonferrous Metals Soc. China 24 (2014) 798–805. [25] Y.V.R.K. Prasad, H.L. Gegel, S.M. Doraivelu, J.C. Malas, J.T. Morgan, K.A. Lark, D.R. Barke, Modeling of dynamic material behavior in hot deformation: forging of Ti-6242, Metall. Trans. A 15 (1984) 1883–1892. [26] T. Zhang, Y. Tao, X. Wang, Constitutive behavior, microstructural evolution and processing map of extruded Al–1.1Mn–0.3Mg–0.25RE alloy during hot compression, Trans. Nonferrous Metals Soc. China 24 (2014) 1337–1345. [27] S.Y. Park, W.J. Kim, Difference in the hot compressive behavior and processing maps between the as-cast and homogenized Al-Zn-Mg-Cu (7075) alloys, J. Mater. Sci. Technol. 32 (2016) 660–670. [28] H. Wu, S.P. Wen, H. Huang, K.Y. Gao, X.L. Wu, Hot deformation behavior and processing map of a new type Al-Zn-Mg-Er-Zr alloy, J. Alloys Compd. 685 (2016) 869–880. [29] J. Huang, G. Gray III, Serrated flow and negative rate sensitivity in Al-Li base alloys, Scr. Metall. Mater. 24 (1990) 85–90. [30] S. Guo, D. Li, X. Wu, X. Xu, P. Du, J. Hu, Characterization of hot deformation behavior of a Zn-10.2 Al-2.1 Cu alloy using processing maps, Mater. Des. 41 (2012) 158–166. [31] X. Yang, L. Chai, W. Huang, Y. Ma, Z. Zhang, EBSD analysis on restoration mechanism of as-extruded AA2099 Al-Li alloy after various thermomechanical processes, Mater. Chem. Phys. 191 (2017) 99–105. [32] X. Wang, S. Bo, D. Lei, J. Li, Effects of original orientation combination on substructure characteristics during continuous dynamic recrystallization in an extruded Al-Cu-Li alloy, Mater. Charact. 130 (2017) 113–122. [33] V. Taupin, J. Chevy, C. Fressengeas, Effects of grain-to-grain interactions on shear strain localization in Al–Cu–Li rolled sheets, Int. J. Solids Struct. 99 (2016) 71–81. [34] K. Huang, R.E. Logé, A review of dynamic recrystallization phenomena in metallic materials, Mater. Des. 111 (2016) 548–574. [35] L. Jiang, M.T. Pérez-Prado, P.A. Gruber, E. Arzt, O.A. Ruano, M.E. Kassner, Texture, microstructure and mechanical properties of equiaxed ultrafine-grained Zr fabricated by accumulative roll bonding, Acta Mater. 56 (2008) 1228–1242. [36] M. Bacca, D.R. Hayhurst, R.M. Mcmeeking, Continuous dynamic recrystallization during severe plastic deformation, Mech. Mater. 90 (2015) 148–156. [37] Q. Contrepois, C. Maurice, J.H. Driver, Hot rolling textures of Al–Cu–Li and Al–Zn–Mg–Cu aeronautical alloys: experiments and simulations to high strains, Mater. Sci. Eng. A 527 (2010) 7305–7312. [38] Q. Zhao, Z. Liu, S. Li, T. Huang, P. Xia, Evolution of the brass texture in an Al-Cu-Mg alloy during hot rolling, J. Alloys Compd. 691 (2017) 786–799. [39] D. Tsivoulas, P.B. Prangnell, The effect of Mn and Zr dispersoid-forming additions on recrystallization resistance in Al–Cu–Li AA2198 sheet, Acta Mater. 77 (2014) 1–16.

Acknowledge Scientific/technical assistance from Microstructure at CSU is acknowledged.

Institute

for

Materials

References [1] Y.E. Ma, Z. Zhao, B. Liu, W. Li, Mechanical properties and fatigue crack growth rates in friction stir welded nugget of 2198-T8 Al–Li alloy joints, Mater. Sci. Eng. A 569 (2013) 41–47. [2] A.I. Karayan, K. Jata, M. Velez, H. Castaneda, On exfoliation corrosion of alloy 2060 T8E30 in an aggressive acid environment, J. Alloys Compd. 657 (2016) 546–558. [3] E. Gumbmann, F. De Geuser, C. Sigli, A. Deschamps, Influence of Mg, Ag and Zn minor solute additions on the precipitation kinetics and strengthening of an Al-CuLi alloy, Acta Mater. 133 (2017) 172–185. [4] B. Decreus, A. Deschamps, F.D. Geuser, P. Donnadieu, C. Sigli, The influence of Cu/ Li ratio on precipitation in Al–Cu–Li–x alloys, Acta Mater. 61 (2013) 2207–2218. [5] B. Wu, M.Q. Li, D.W. Ma, The flow behavior and constitutive equations in isothermal compression of 7050 aluminum alloy, Mater. Sci. Eng. A 542 (2012) 79–87. [6] Y.C. Lin, Q. Li, Y. Xia, L. Li, A phenomenological constitutive model for high temperature flow stress prediction of Al–Cu–Mg alloy, Mater. Sci. Eng. A 534 (2012) 654–662. [7] Y. Liu, C. Geng, Q. Lin, Y. Xiao, J. Xu, W. Kang, Study on hot deformation behavior and intrinsic workability of 6063 aluminum alloys using 3D processing map, J. Alloys Compd. 713 (2017) 212–221. [8] B. Li, Q. Pan, Z. Yin, Microstructural evolution and constitutive relationship of Al–Zn–Mg alloy containing small amount of Sc and Zr during hot deformation based on Arrhenius-type and artificial neural network models, J. Alloys Compd. 584 (2014) 406–416. [9] J. Li, Y. Lio, Y. Wang, B. Liu, B. Lu, X. Liang, Constitutive equation and processing map for hot compressed as-cast Ti-43Al-4Nb-1.4W-0.6B alloy, Trans. Nonferrous Metals Soc. China 23 (2013) 3383–3391. [10] K. Ihara, Y. Miura, Dynamic recrystallization in Al–Mg–Sc alloys, Mater. Sci. Eng. A 387–389 (2004) 647–650. [11] M. Zhou, Y.C. Lin, J. Deng, Y. Jiang, Hot tensile deformation behaviors and constitutive model of an Al–Zn–Mg–Cu alloy, Mater. Des. 59 (2014) 141–150. [12] F. Zhang, J.L. Sun, J. Shen, X.D. Yan, J. Chen, Flow behavior and processing maps of 2099 alloy, Mater. Sci. Eng. A 613 (2014) 141–147. [13] S.C. Hogg, I.G. Palmer, L.G. Thomas, P.S. Grant, Processing, microstructure and property aspects of a spraycast Al–Mg–Li–Zr alloy, Acta Mater. 55 (2007) 1885–1894. [14] Q. Yang, X. Wang, X. Li, Z. Deng, Z. Jia, Z. Zhang, G. Huang, Q. Liu, Hot deformation behavior and microstructure of AA2195 alloy under plane strain compression, Mater. Charact. 131 (2017) 500–507. [15] R. Zhu, Q. Liu, J. Li, S. Xiang, Y. Chen, X. Zhang, Dynamic restoration mechanism and physically based constitutive model of 2050 Al–Li alloy during hot compression, J. Alloys Compd. 650 (2015) 75–85. [16] N. Nayan, S. Murty, S. Chhangani, A. Prakash, M. Prasad, Effect of temperature and strain rate on hot deformation behavior and microstructure of Al-Cu-Li alloy, J. Alloys Compd. 723 (2017) 548–558. [17] H. Yin, H. Li, X. Su, D. Huang, Processing maps and microstructural evolution of

64