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Study on the properties and microstructure of dispersion strengthened copper alloy deformed at high temperatures
Journal of Alloys and Compounds 479 (2009) 401–408
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Study on the properties and microstructure of dispersion strengthened copper alloy deformed at high temperatures K. Shen ∗ , M.P. Wang, S.M. Li School of Materials Science and Engineering, Central South University, Changsha, 410083, PR China
a r t i c l e
i n f o
Article history: Received 6 November 2008 Received in revised form 9 December 2008 Accepted 15 December 2008 Available online 25 December 2008 Keywords: Metals Powder metallurgy Equation of state Grain boundaries Strain
a b s t r a c t The changes of flow stress, hardness and microstructure of the Cu–0.23vol%Al2 O3 alloy in the high temperature plastic deformation process have been investigated. The differences of flow stress of this alloy compressed under different conditions are significant, yet, mainly including three different stages. The constitutive equation for its peak yield stress, strain rate and temperature was also established, which can evaluate its peak yield stress better. With the increasing of temperature and strain rate, the changes of hardness and microstructure also have its own uniqueness. The higher strain rate is favorable to the formation of fine subgrains with a clear subgrain boundary. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Dispersion strengthened copper alloys possess excellent high strength and high values of electrical and thermal conductivities [1,2]. And they have been widely used as contacts, lead wires, electrodes, vacuum technique parts and conductors in high temperature electrical applications [1,3,4]. Therefore, up to now, many preparation methods have been developed to prepare this kind of material [5–8], such as internal oxidation, mechanical alloying, in situ and ex situ spray deposition. Cu–Al2 O3 alloy prepared by internal oxidation has been investigated for many years, yet, these investigations were mainly focused on the preparation methods, room deformation and annealing behaviors. Through investigation, it is well known that the excellent properties of Cu–Al2 O3 alloys are mainly attributed to the presence of thermally stable ␥alumina particles dispersed in the copper matrix. Such dispersed particles can increase the threshold stress for dislocation glide and cause the generation of additional dislocations around particles during plastic deformation. These effects lead to increases in the strength of the Cu–Al2 O3 alloys [9]. When an Cu–Al2 O3 alloy strip is cold rolled, a dislocation cell substructure at small or moderate strains, or a band substructure at large strains, develops [10]. The size of the cells or bands, or the band boundary misorientation, depends for a given Cu–Al2 O3 alloy on its deformation characteristics, mechanical properties and annealing behavior. Moreover, the
∗ Corresponding author. Tel.: +86 731 8830264. E-mail address: [email protected] (K. Shen). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.12.060
fine dispersion alumina particles play a role as barriers to the movement of subgrain or grain boundaries, which retards the recovery and recrystallization of the Cu–Al2 O3 alloy, even at high temperatures [11]. However, with the development of high and new technologies, the requirement for properties of Cu–Al2 O3 alloys are gradually increased, especially for the high temperature mechanical properties, yet, the investigation about this is very few [12–15]. Therefore, it is very useful and necessary to research the properties and microstructure of Cu–Al2 O3 alloy deformed at high temperatures. 2. Experimental procedure Experimental material used in this study was Cu–0.23vol%Al2 O3 alloy produced by a simplified internal oxidation method. The internal oxidation procedure consisted of: induction melting of Cu–Al alloy → nitrogen atomization → mixing of the atomized Cu–Al alloy with an oxidant → internal oxidation at 1273 K for 1 h → hydrogen reduction at 1173 K for 1 h → vacuum hot-processing at 1223 K for 3 h under a pressure of 27 MPa and a pressure of 1.33 × 10−2 Pa → hot extrusion at 1203 K (extrusion rate 50:1) → cold drawing to the bar with the diameter of ˚16. Hot compression test was performed on the Gleeble-1500 thermal simulator with protection of argon atmosphere. The cutting way and shape of hot compressed samples (˚8(±0.05) × 12(±0.1)) are shown in Fig. 1. The sample was first heated to the setting temperature with a rate of 5 ◦ C/s, and kept it for 3 min, then the sample was hot compressed to the setting deformation amount with a constant temperature and a strain rate, at last, in order to get the high temperature deformation microstructure, the compressed sample was quickly water-cooled to the room temperature. In addition, one lubricant (graphite75% + 46# engine oil 20% + nitric acid three toluene fat 5%) was used to reduce the friction and uneven deformation. The hot compression conditions were as follows: deformation temperature: 350, 500, 650, 800, and 950 ◦ C; strain rate: 0.001, 0.01, 0.1, 1 and 20 s−1 ; deformation amount: 80%. The variations of stress and strain were monitored continuously by a personal computer equipped with an automatic data acquisition system. The true stress and
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Fig. 1. Cutting way and shape of thermal simulation samples (a) cutting way and (b) sample shape.
true strain were derived from the measurement of displacement and load according to the following relationships:
εtr = ln
l0 l0 − l
F(l0 − l) tr = S0 l0
3. Results and discussion 3.1. Effect of compression conditions on stress–strain curves
(1)
(2)
where l0 is the original length of sample, l is the displacement, F is the load, S0 is the original cross-sectional area of sample. The hardness of specimens under different compression conditions was characterized by the Vickers hardness under a load of 3 kg for 30 s. Grain morphology and other intrinsic microstructure features of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions were revealed using an etchant mixture of 5 ml ferric chloride, 25 ml hydrochloric acid and 100 ml distilled water. The polished and etched specimens were examined in a Leica DMILM metallographic microscope. The transmission electron microscopy (TEM) observation of the Cu–0.23vol%Al2 O3 alloys was performed on a transmission electron microscope FEI TECNAIG2 at 200 kV. Thin foil samples were electropolished in a twin-jet instrument using a mixture of 30% hydrogen nitrate and alcohol as balance.
The effect of temperature on the true stress–true strain curves of the Cu–0.23vol%Al2 O3 alloy is shown in Fig. 2. From it, we can see that, the change regularity of flow stress of this alloy compressed under different conditions is significantly different. The curves of true stress–true strain can be mainly characterized by an initial sharp increase stage, steady stage, and flow stress softening stage (as shown in Fig. 2(d)). For the initial sharp increase stage, if the strain rate is the same, peak yield stress of this alloy decreases with the increasing of deformation temperature, yet, when the compression temperature is the same, its peak yield stress increases with the increasing of strain rate. For the secondary stage, from Fig. 2, we can see that the higher deformation temperature is, the easier does steady flow stress stage appear, and the smaller is the differential value between peak yield stress and steady flow stress, which
Fig. 2. The effect of temperature on the true stress-true strain curves of low concentration Cu–0.23vol%Al2 O3 alloy (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.01 s−1 , (c) ε˙ = 0.1 s−1 , (d) ε˙ = 1 s−1 and (e) ε˙ = 20 s−1
K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
indicates the dynamic balance between working hardening and working softening can be reached in a shorter time. However, when the strain rate is increased to 20 s−1 , the fluctuation phenomenon of flow stress can be first observed before the appearance of steady stage in the lower temperature compression, and this kind of fluctuation may be caused by the fact that work softening resulted from dynamic recovery and dynamic recrystallization cannot get a balance with work hardening in time. However, when the temperature is increased to above 800 ◦ C, the rate of softening resulted from dynamic recovery or dynamic recrystallization may be much higher enough, and can get a balance with the rate of work hardening, therefore, the fluctuation phenomenon of flow stress disappears at or above 800 ◦ C (Fig. 2(e)). For the flow stress softening stage following the steady stage, as mentioned above, it can be observed mainly during the lower temperature deformation (below 650 ◦ C). Which may be resulted from both dislocations reacting with each other and increasing of subgrain size, and because it is difficult to move for dislocation lines during low temperature deformation, the density of dislocation in front of particles will be increased quickly, at last, it is very easy to occur for the accumulated unlike dislocation reacting with each other [7], and form a clear subgrain boundary (as mentioned in the following parts). 3.2. Effect of compression on the peak yield stress From the true stress–true strain curves of the Cu–0.23vol%Al2 O3 alloy, we can find that the interrelation between flow stress, strain rate and temperature (at constant strain) is close, therefore, in order to better understand the high temperature plastic deformation behaviors of this alloy, it is very necessary to investigate their relationship. It is well known that if the high temperature deformation
403
of metal materials is a thermal activation process, the interrelation between flow stress, strain rate and temperature can be analyzed by using a phenomenological approach comprising the equation proposed by Sellars and Tegart [8,9]:
ε˙ = AF() exp −
Q (RT )
(3)
where ε˙ is strain rate, A is constant, Q is activation energy, R is gas constant, T is temperature, F() is the function of flow stress, and it can be expressed as F() = [sinh(˛)]n
(4)
If ˛ > 1.2, F() = exp(ˇ), and if ˛ < 0.8, F() = n , where ˛, ˇ and n are constants, and they satisfy the relationship ˛ = ˇ/n. According to the Eq. (3), we can get the following two equations: ε˙ = B exp(ˇ)
(5)
ε˙ = B n
(6)
where B and B are constant, respectively. And the following two equations can be obtained: ln ε˙ = − ln B + ˇ
(7)
ln ε˙ = − ln B + n ln
(8)
According to Eqs. (3) and (4), the following equation also can be obtained
Q
ε˙ = A[sinh(˛)]n exp −
RT
(9)
Fig. 3. The relationship between peak value stress and other variable parameters (a) ln ε˙ − , (b) ln ε˙ − ln , (c) ln ε˙ − ln[sinh(˛)], (d) ln[sinh(˛)] − T−1 and (e) ln Z − ln[sinh(˛)].
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Table 1 Material parameters for establishing constitutive equation of Cu–Al2 O3 alloy. ˇ-Value 54.9165 × 10
Cu–Al2 O3 alloy
˛-Value
n-Value −3
−3
12.484 × 10
4.39909
ln ε˙ = n ln[sinh(˛)] + −
RT
+ ln A
350 ◦ C
(10)
From the Eqs. (7), (8) and (10), we know that ln ε˙ and , ln ε˙ and ln , ln ε˙ and ln[sinh(˛)] all satisfy the linear relationship. The relationship curves of ln ε˙ − , ln ε˙ − ln and ln ε˙ − ln[sinh(˛)] for different deformation conditions are shown in Fig. 3. From it, we can find that all of them basically satisfy the linear relationship, and the material parameters obtained from them are shown in Table 1, which indicates that peak yield stress and strain rate satisfy the hyperbolic sine relationship very well, and the high temperature compression of the Cu–0.23vol%Al2 O3 alloy is also a thermal activation process. Because the deformation temperature is also very important to the deformation process, it is very necessary to analyze the quantitative relationship between temperature and peak yield stress. From the Eq. (10), we can see that if the activation energy Q will not change with the increasing of temperature, the relationship between the term ln[sinh(˛)] and T−1 should be linear relationship. This result is proved by the experimental data as shown in Fig. 3(d), yet, those activation energy values compressed at 950 ◦ C greatly deviate from the fitted line (as shown in the elliptic part in Fig. 3(d)). This may be resulted from the formation of crack on the boundary of alumina particles caused by the different coefficient of thermal expansion between alumina particle and copper matrix. According the slope rate of the fitted lines (=Q/nR value), we can obtain the average activation energy Q (as shown in Table 1). Compared with the creep activation energy (253.3 kJ/mol) of Glidcop Al-15 alloy [10], it is much lower. This may be resulted from the different concentration of alumina particle and different deformation ways. In order to establish a constitutive equation between peak yield stress, strain rate and temperature for the Cu–0.23vol%Al2 O3 alloy, the Zener–Hollomon parameter also need to be introduced, just as following: Z = ε˙ exp
ln A value
99.848
11.65218
Table 2 Comparison between measured and calculated peak value yield stress value.
Then we can get the following equation
Q
Q value (kJ/mol)
Q
(11)
RT
That is, Z = A sinh [˛]n
(12)
To take the logarithm of Eq. (12), we can get ln Z = ln A + n ln[sinh(˛)]
(13)
According to the Eq. (13), we know that both n and ln A values can be obtained from the relationship curve of ln Z − ln[sinh(˛)] (as shown in Fig. 3(e)). Table 1 shows the measured values of n and ln A. At last, the constitutive equation of the Cu–0.23vol%Al2 O3 alloy
0.001 s−1 0.01 s−1 0.1 s−1 1 s−1 20 s−1
500 ◦ C
650 ◦ C
800 ◦ C
950 ◦ C
M
C
M
C
M
C
M
C
M
C
93 136 183 214 263
92 136 181 214 264
34 69 119 153 200
34 70 117 152 200
– 46 74 112 162
– 43 71 110 160
– 32 56 92 121
– 29 53 90 120
– 18 29 62 89
– 16 27 60 88
M, measured value and C, calculated value.
compressed at high temperatures can be established as follows:
ε˙ = [sinh(0.0124836)]
4.39909
exp
11.65218 − 99.848 × 103 RT
(14) Using this constitutive equation, we can calculate the theoretical peak yield stress of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions, the measured and calculated values are shown in Table 2. From it, we can see that the calculated values are very close to the measured values. Therefore, this constitutive equation can be used to evaluate the peak yield stress of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions better. 3.3. Effect of compression conditions on the hardness Table 3 shows the hardness values of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions. And the corresponding relationship curves are shown in Fig. 4. From Fig. 4(a), we can see that, the hardness value of the Cu–0.23vol%Al2 O3 alloy compressed at the same temperature is first increased with the increasing of strain rate, then after reaching the peak value it begin to decrease with further increasing of strain rate. Taking a careful analysis, we can find that those peak hardness values compressed at or under 500 ◦ C all appear in the strain rate of 0.1 s−1 , yet, when the compression temperature is in the range of 650–950 ◦ C, their peak hardness values appear in the strain rate of 1 s−1 . This difference is mainly resulted from the interaction between work hardening and dynamic recovery or dynamic recrystallization. We know that, during the same temperature compression, with the increasing of strain rate, on the one hand, proliferation rate of dislocation will be speeded up, leading to the increase in the rate of work hardening for this alloy; on the other hand, due to the pinning effect of dispersion alumina particles, a large number of dislocation lines will be stopped by the alumina particles and seriously react with each other, and leading to the increasing in the rate of dynamic recovery. At last, the interaction between work hardening and working soft-
Table 3 The hardness value of Cu–0.23vol%Al2 O3 alloy compressed under different conditions. Compression conditions
Cu–0.23vol%Al2 O3 alloy
ε˙
0.001 0.01 0.1 1 20
The hardness value (kgf/mm2 ) Original
350 ◦ C
500 ◦ C
650 ◦ C
800 ◦ C
950 ◦ C
120.6
114.9 126.2 127.9 125.2 124.2
100 109.7 116.9 110.7 115.4
96.3 97 107.6 114.8 112.1
72.8 94.6 94.8 103 93
– 83.8 86.5 97.5 91.2
K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
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Fig. 4. Effect of compression conditions on the hardness of Cu–0.23vol%Al2 O3 alloy. (a) Effect of strain rate on the hardness. (b) Effect of temperature on the hardness.
Fig. 5. The original structure of low concentration Cu–0.23vol%Al2 O3 alloy used for thermal simulation. (a) Longitudinal direction and (b) transverse direction.
ening will inevitably leads to the appearance of peak hardness value in a certain strain rate. In addition, as can be seen from Fig. 4(b), with the increasing of compression temperature, the occurrence of dynamic recovery or dynamic recrystallization becomes more and more significant, therefore, a downward trend of hardness for the Cu–0.23vol%Al2 O3 alloy compressed in the same strain rate can be observed.
3.4. Effect of compression conditions on the microstructure In order to explain the change characteristics of the true stress–true strain curves and hardness of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions, it is quite necessary to study the effect of compression conditions on its microstructure. The original microstructure of the Cu–0.23vol%Al2 O3 alloy is
Fig. 6. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 350 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
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K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
Fig. 7. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
Fig. 8. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 650 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
shown in Fig. 5. Highly elongated fibre structure can be observed in the longitudinal direction, and these fibres are arranged to the drawing direction (Fig. 5(a)). Yet, a granular microstructure appears on the transverse cross-section (Fig. 5(b)). The metallographic microstructures of Cu–0.23vol%Al2 O3 alloy compressed under different conditions are shown in Figs. 6–9. From Fig. 6, it can be seen that, the original fibre microstructure is greatly weakened, especially for the case of strain rate 0.001 s−1 yet, dynamic recrystallization grains cannot be observed. When the compression temperature is increased to 500 ◦ C, a large number of fine
dynamic recrystallization grains are formed after compression by three strain rates (as shown in Fig. 7). The most dynamic recrystallization grains can be observed on the overall matrix for the case of strain rate 0.001 s−1 (Fig. 7(a)). With the increasing of strain rate, its deformation becomes more and more non-homogeneous, and length–width ratio of fibre is also increased (Fig. 7(c)). In addition, though fibre structure still can be seen for the case of the higher strain rate (20 s−1 ), yet, most of those fibre structure have been penetrated by the dynamic recrystallization grains thoroughly (Fig. 7(c)). When the compression temperature is increased
Fig. 9. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 800 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
407
Fig. 10. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
Fig. 11. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 800 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
to 650 ◦ C, the size of dynamic recrystallization grains is significantly increased (Fig. 8(a)), and the original fibre structure has been divided into the smaller grain cells by dynamic recrystallization grains (as arrow directions in Fig. 8(b) and (c)). When the temperature is further increased to 800 ◦ C, from Fig. 9, we can see that the size of dynamic recrystallization grain is also increased greatly, especially for the case of strain rate 0.001 s−1 . For the different high temperature compression process, both the interaction between dislocation and dispersion particles, and subgrain structure are different for the Cu–0.23vol%Al2 O3 alloy. Therefore, there is need to analyze their TEM microstructure under different conditions. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C with different strain rates is shown in Fig. 10. From it, we can see that, with the increasing of strain rate, the size of subgrain with an equiaxed shape is gradually decreased. When the sample is compressed in the strain rate 0.001 s−1 , subgrain size is the largest (about 2–3 m), and the density of dislocation is much lower, only those dispersion particles distributed in some coarser subgrains are tangled by a few dislocation lines (Fig. 10(a)). With the increasing of strain rate to 0.1 s−1 , a reduced subgrain size can be seen (about 1.5–2 m), yet, the density of dislocation is much higher than that of in strain rate 0.001 s−1 , and most of these dislocation lines mainly distribute around the subgrain boundary. Besides, both the combination of some subgrains, and transformation of small-angle to wide-angle grain boundary can be observed in Fig. 10(b). With further increasing of strain rate (to 20 s−1 ), it can be see that the subgrain size is reduced to only about 0.5–1 m at this moment, and the clear
subgrain boundary can be also observed, yet, the density of dislocation is lower than that of in strain rate 0.1 s−1 , which suggests the higher strain rate is favorable to the formation of subgrain structure. When the temperature is increased to 800 ◦ C, the size of subgrain is much larger than that of compressed at 500 ◦ C in the same strain rate, and the higher strain rate is also favorable to the formation of clear subgrain boundary (Fig. 11), which is the same as that of the lower temperature compression of the Cu–0.23vol%Al2 O3 alloy. 4. Conclusions The properties and microstructure of the Cu–0.23vol%Al2 O3 alloy compressed at high temperatures have been investigated. The results are as follows: (1) The curves of true stress–true strain for the Cu–0.23vol%Al2 O3 alloy compressed at high temperatures can be mainly characterized by an initial sharp increase stage, steady stage, and flow stress softening stage. Both the peak yield stress and strain rate satisfies the hyperbolic sine relationship, indicating its high temperature compression being a thermal activation process. Based on the calculated material parameters, the constitutive equation for peak yield stress, strain rate and temperature was established as follows: ε˙ = [sinh(0.0124836)]4.39909 exp(11.65218 − 99.848 × 103 /RT ). (2) The hardness of the Cu–0.23vol%Al2 O3 alloy compressed at the same temperature is first increased with the increasing of strain
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rate, then after reaching the peak value it begins to decrease with further increasing of strain rate. With the increasing of compression temperature, a downward trend of hardness for this alloy compressed in the same strain rate can be observed. (3) With the increasing of compression temperature, the size and number of dynamic recrystallization grains in the alloy matrix are increased. Yet, when deformed at the same temperature, with the increasing of strain rates, uneven distribution of metallographical microstructure is also strengthened, and the size of subgrain is gradually decreased; yet, the dislocation density is increased at first, and then followed by a decreasing. The higher strain rate is favorable to the formation of fine subgrains with a clear subgrain boundary. Acknowledgement This study was supported by the National High-Tech Research and Development program of China (863 Program).
References [1] A.V. Nadkarni, Metall. Soc. AIME (1984) 77. [2] J.E. Synk, K. Vendula, Mater. Sci. Technol. 3 (1987) 72–75. [3] J.S. Lee, Y.C. Kim, S.H. Lee, S.H. Ahn, N.J. Kim, Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 35 (2004) 493–502. [4] S. Sato, T. Hatano, T. Kuroda, K. Furoya, S. Hara, M. Enoed, et al., J. Nucl. Mater. 263 (1998) 265–270. [5] J.E. Perez, D.G. Morris, Scr. Met. Mater. 31 (1994) 231–235. [6] S. Ziyuan, W. Deqing, J. Mater. Sci. Lett. 17 (1998) 477–479. [7] V.M. Rajkovic, M.V. Mitkov, Int. J. Powder Metall. 36 (2000) 45–49. [8] M.X. Guo, M.P. Wang, K. Shen, L.F. Cao, Z. Li, Z. Zhang, J. Alloys Compd. 460 (2008) 585–589. [9] J.W. Martin, Micromechanisms in Particle-hardened Alloys, Cambridge University Press, Cambridge, 1980, p. 50. [10] I. Baker, J.W. Martin, Met. Sci. 17 (1983) 459–467. [11] J.L. Brimhall, M.J. Klein, R.A. Huggins, Acta Metall. 14 (1966) 45–66. ˜ [12] R.G. Espinoza, R.H. Palma, A.O. Sepúlveda, A. Zúniga, Mater. Sci. Eng. A 498 (2008) 397–403. [13] R.H. Palma, A. Sepúlveda, R. Espinoza, Jesús DiánezF M., José M. Criado, M. Jesús Sayagués, Mater. Sci. Eng. A 391 (2005) 60–65. [14] T.S. Srivatsan, N. Narendra, J.D. Troxell, Mater. Des. 21 (2000) 191–198. [15] T.S. Srivatsan, K. Dhana Singh, J.D. Troxell, Mater. Lett. 28 (1996) 423–429.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Study on the properties and microstructure of dispersion strengthened copper alloy deformed at high temperatures K. Shen ∗ , M.P. Wang, S.M. Li School of Materials Science and Engineering, Central South University, Changsha, 410083, PR China
a r t i c l e
i n f o
Article history: Received 6 November 2008 Received in revised form 9 December 2008 Accepted 15 December 2008 Available online 25 December 2008 Keywords: Metals Powder metallurgy Equation of state Grain boundaries Strain
a b s t r a c t The changes of flow stress, hardness and microstructure of the Cu–0.23vol%Al2 O3 alloy in the high temperature plastic deformation process have been investigated. The differences of flow stress of this alloy compressed under different conditions are significant, yet, mainly including three different stages. The constitutive equation for its peak yield stress, strain rate and temperature was also established, which can evaluate its peak yield stress better. With the increasing of temperature and strain rate, the changes of hardness and microstructure also have its own uniqueness. The higher strain rate is favorable to the formation of fine subgrains with a clear subgrain boundary. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Dispersion strengthened copper alloys possess excellent high strength and high values of electrical and thermal conductivities [1,2]. And they have been widely used as contacts, lead wires, electrodes, vacuum technique parts and conductors in high temperature electrical applications [1,3,4]. Therefore, up to now, many preparation methods have been developed to prepare this kind of material [5–8], such as internal oxidation, mechanical alloying, in situ and ex situ spray deposition. Cu–Al2 O3 alloy prepared by internal oxidation has been investigated for many years, yet, these investigations were mainly focused on the preparation methods, room deformation and annealing behaviors. Through investigation, it is well known that the excellent properties of Cu–Al2 O3 alloys are mainly attributed to the presence of thermally stable ␥alumina particles dispersed in the copper matrix. Such dispersed particles can increase the threshold stress for dislocation glide and cause the generation of additional dislocations around particles during plastic deformation. These effects lead to increases in the strength of the Cu–Al2 O3 alloys [9]. When an Cu–Al2 O3 alloy strip is cold rolled, a dislocation cell substructure at small or moderate strains, or a band substructure at large strains, develops [10]. The size of the cells or bands, or the band boundary misorientation, depends for a given Cu–Al2 O3 alloy on its deformation characteristics, mechanical properties and annealing behavior. Moreover, the
∗ Corresponding author. Tel.: +86 731 8830264. E-mail address: [email protected] (K. Shen). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.12.060
fine dispersion alumina particles play a role as barriers to the movement of subgrain or grain boundaries, which retards the recovery and recrystallization of the Cu–Al2 O3 alloy, even at high temperatures [11]. However, with the development of high and new technologies, the requirement for properties of Cu–Al2 O3 alloys are gradually increased, especially for the high temperature mechanical properties, yet, the investigation about this is very few [12–15]. Therefore, it is very useful and necessary to research the properties and microstructure of Cu–Al2 O3 alloy deformed at high temperatures. 2. Experimental procedure Experimental material used in this study was Cu–0.23vol%Al2 O3 alloy produced by a simplified internal oxidation method. The internal oxidation procedure consisted of: induction melting of Cu–Al alloy → nitrogen atomization → mixing of the atomized Cu–Al alloy with an oxidant → internal oxidation at 1273 K for 1 h → hydrogen reduction at 1173 K for 1 h → vacuum hot-processing at 1223 K for 3 h under a pressure of 27 MPa and a pressure of 1.33 × 10−2 Pa → hot extrusion at 1203 K (extrusion rate 50:1) → cold drawing to the bar with the diameter of ˚16. Hot compression test was performed on the Gleeble-1500 thermal simulator with protection of argon atmosphere. The cutting way and shape of hot compressed samples (˚8(±0.05) × 12(±0.1)) are shown in Fig. 1. The sample was first heated to the setting temperature with a rate of 5 ◦ C/s, and kept it for 3 min, then the sample was hot compressed to the setting deformation amount with a constant temperature and a strain rate, at last, in order to get the high temperature deformation microstructure, the compressed sample was quickly water-cooled to the room temperature. In addition, one lubricant (graphite75% + 46# engine oil 20% + nitric acid three toluene fat 5%) was used to reduce the friction and uneven deformation. The hot compression conditions were as follows: deformation temperature: 350, 500, 650, 800, and 950 ◦ C; strain rate: 0.001, 0.01, 0.1, 1 and 20 s−1 ; deformation amount: 80%. The variations of stress and strain were monitored continuously by a personal computer equipped with an automatic data acquisition system. The true stress and
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K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
Fig. 1. Cutting way and shape of thermal simulation samples (a) cutting way and (b) sample shape.
true strain were derived from the measurement of displacement and load according to the following relationships:
εtr = ln
l0 l0 − l
F(l0 − l) tr = S0 l0
3. Results and discussion 3.1. Effect of compression conditions on stress–strain curves
(1)
(2)
where l0 is the original length of sample, l is the displacement, F is the load, S0 is the original cross-sectional area of sample. The hardness of specimens under different compression conditions was characterized by the Vickers hardness under a load of 3 kg for 30 s. Grain morphology and other intrinsic microstructure features of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions were revealed using an etchant mixture of 5 ml ferric chloride, 25 ml hydrochloric acid and 100 ml distilled water. The polished and etched specimens were examined in a Leica DMILM metallographic microscope. The transmission electron microscopy (TEM) observation of the Cu–0.23vol%Al2 O3 alloys was performed on a transmission electron microscope FEI TECNAIG2 at 200 kV. Thin foil samples were electropolished in a twin-jet instrument using a mixture of 30% hydrogen nitrate and alcohol as balance.
The effect of temperature on the true stress–true strain curves of the Cu–0.23vol%Al2 O3 alloy is shown in Fig. 2. From it, we can see that, the change regularity of flow stress of this alloy compressed under different conditions is significantly different. The curves of true stress–true strain can be mainly characterized by an initial sharp increase stage, steady stage, and flow stress softening stage (as shown in Fig. 2(d)). For the initial sharp increase stage, if the strain rate is the same, peak yield stress of this alloy decreases with the increasing of deformation temperature, yet, when the compression temperature is the same, its peak yield stress increases with the increasing of strain rate. For the secondary stage, from Fig. 2, we can see that the higher deformation temperature is, the easier does steady flow stress stage appear, and the smaller is the differential value between peak yield stress and steady flow stress, which
Fig. 2. The effect of temperature on the true stress-true strain curves of low concentration Cu–0.23vol%Al2 O3 alloy (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.01 s−1 , (c) ε˙ = 0.1 s−1 , (d) ε˙ = 1 s−1 and (e) ε˙ = 20 s−1
K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
indicates the dynamic balance between working hardening and working softening can be reached in a shorter time. However, when the strain rate is increased to 20 s−1 , the fluctuation phenomenon of flow stress can be first observed before the appearance of steady stage in the lower temperature compression, and this kind of fluctuation may be caused by the fact that work softening resulted from dynamic recovery and dynamic recrystallization cannot get a balance with work hardening in time. However, when the temperature is increased to above 800 ◦ C, the rate of softening resulted from dynamic recovery or dynamic recrystallization may be much higher enough, and can get a balance with the rate of work hardening, therefore, the fluctuation phenomenon of flow stress disappears at or above 800 ◦ C (Fig. 2(e)). For the flow stress softening stage following the steady stage, as mentioned above, it can be observed mainly during the lower temperature deformation (below 650 ◦ C). Which may be resulted from both dislocations reacting with each other and increasing of subgrain size, and because it is difficult to move for dislocation lines during low temperature deformation, the density of dislocation in front of particles will be increased quickly, at last, it is very easy to occur for the accumulated unlike dislocation reacting with each other [7], and form a clear subgrain boundary (as mentioned in the following parts). 3.2. Effect of compression on the peak yield stress From the true stress–true strain curves of the Cu–0.23vol%Al2 O3 alloy, we can find that the interrelation between flow stress, strain rate and temperature (at constant strain) is close, therefore, in order to better understand the high temperature plastic deformation behaviors of this alloy, it is very necessary to investigate their relationship. It is well known that if the high temperature deformation
403
of metal materials is a thermal activation process, the interrelation between flow stress, strain rate and temperature can be analyzed by using a phenomenological approach comprising the equation proposed by Sellars and Tegart [8,9]:
ε˙ = AF() exp −
Q (RT )
(3)
where ε˙ is strain rate, A is constant, Q is activation energy, R is gas constant, T is temperature, F() is the function of flow stress, and it can be expressed as F() = [sinh(˛)]n
(4)
If ˛ > 1.2, F() = exp(ˇ), and if ˛ < 0.8, F() = n , where ˛, ˇ and n are constants, and they satisfy the relationship ˛ = ˇ/n. According to the Eq. (3), we can get the following two equations: ε˙ = B exp(ˇ)
(5)
ε˙ = B n
(6)
where B and B are constant, respectively. And the following two equations can be obtained: ln ε˙ = − ln B + ˇ
(7)
ln ε˙ = − ln B + n ln
(8)
According to Eqs. (3) and (4), the following equation also can be obtained
Q
ε˙ = A[sinh(˛)]n exp −
RT
(9)
Fig. 3. The relationship between peak value stress and other variable parameters (a) ln ε˙ − , (b) ln ε˙ − ln , (c) ln ε˙ − ln[sinh(˛)], (d) ln[sinh(˛)] − T−1 and (e) ln Z − ln[sinh(˛)].
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Table 1 Material parameters for establishing constitutive equation of Cu–Al2 O3 alloy. ˇ-Value 54.9165 × 10
Cu–Al2 O3 alloy
˛-Value
n-Value −3
−3
12.484 × 10
4.39909
ln ε˙ = n ln[sinh(˛)] + −
RT
+ ln A
350 ◦ C
(10)
From the Eqs. (7), (8) and (10), we know that ln ε˙ and , ln ε˙ and ln , ln ε˙ and ln[sinh(˛)] all satisfy the linear relationship. The relationship curves of ln ε˙ − , ln ε˙ − ln and ln ε˙ − ln[sinh(˛)] for different deformation conditions are shown in Fig. 3. From it, we can find that all of them basically satisfy the linear relationship, and the material parameters obtained from them are shown in Table 1, which indicates that peak yield stress and strain rate satisfy the hyperbolic sine relationship very well, and the high temperature compression of the Cu–0.23vol%Al2 O3 alloy is also a thermal activation process. Because the deformation temperature is also very important to the deformation process, it is very necessary to analyze the quantitative relationship between temperature and peak yield stress. From the Eq. (10), we can see that if the activation energy Q will not change with the increasing of temperature, the relationship between the term ln[sinh(˛)] and T−1 should be linear relationship. This result is proved by the experimental data as shown in Fig. 3(d), yet, those activation energy values compressed at 950 ◦ C greatly deviate from the fitted line (as shown in the elliptic part in Fig. 3(d)). This may be resulted from the formation of crack on the boundary of alumina particles caused by the different coefficient of thermal expansion between alumina particle and copper matrix. According the slope rate of the fitted lines (=Q/nR value), we can obtain the average activation energy Q (as shown in Table 1). Compared with the creep activation energy (253.3 kJ/mol) of Glidcop Al-15 alloy [10], it is much lower. This may be resulted from the different concentration of alumina particle and different deformation ways. In order to establish a constitutive equation between peak yield stress, strain rate and temperature for the Cu–0.23vol%Al2 O3 alloy, the Zener–Hollomon parameter also need to be introduced, just as following: Z = ε˙ exp
ln A value
99.848
11.65218
Table 2 Comparison between measured and calculated peak value yield stress value.
Then we can get the following equation
Q
Q value (kJ/mol)
Q
(11)
RT
That is, Z = A sinh [˛]n
(12)
To take the logarithm of Eq. (12), we can get ln Z = ln A + n ln[sinh(˛)]
(13)
According to the Eq. (13), we know that both n and ln A values can be obtained from the relationship curve of ln Z − ln[sinh(˛)] (as shown in Fig. 3(e)). Table 1 shows the measured values of n and ln A. At last, the constitutive equation of the Cu–0.23vol%Al2 O3 alloy
0.001 s−1 0.01 s−1 0.1 s−1 1 s−1 20 s−1
500 ◦ C
650 ◦ C
800 ◦ C
950 ◦ C
M
C
M
C
M
C
M
C
M
C
93 136 183 214 263
92 136 181 214 264
34 69 119 153 200
34 70 117 152 200
– 46 74 112 162
– 43 71 110 160
– 32 56 92 121
– 29 53 90 120
– 18 29 62 89
– 16 27 60 88
M, measured value and C, calculated value.
compressed at high temperatures can be established as follows:
ε˙ = [sinh(0.0124836)]
4.39909
exp
11.65218 − 99.848 × 103 RT
(14) Using this constitutive equation, we can calculate the theoretical peak yield stress of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions, the measured and calculated values are shown in Table 2. From it, we can see that the calculated values are very close to the measured values. Therefore, this constitutive equation can be used to evaluate the peak yield stress of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions better. 3.3. Effect of compression conditions on the hardness Table 3 shows the hardness values of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions. And the corresponding relationship curves are shown in Fig. 4. From Fig. 4(a), we can see that, the hardness value of the Cu–0.23vol%Al2 O3 alloy compressed at the same temperature is first increased with the increasing of strain rate, then after reaching the peak value it begin to decrease with further increasing of strain rate. Taking a careful analysis, we can find that those peak hardness values compressed at or under 500 ◦ C all appear in the strain rate of 0.1 s−1 , yet, when the compression temperature is in the range of 650–950 ◦ C, their peak hardness values appear in the strain rate of 1 s−1 . This difference is mainly resulted from the interaction between work hardening and dynamic recovery or dynamic recrystallization. We know that, during the same temperature compression, with the increasing of strain rate, on the one hand, proliferation rate of dislocation will be speeded up, leading to the increase in the rate of work hardening for this alloy; on the other hand, due to the pinning effect of dispersion alumina particles, a large number of dislocation lines will be stopped by the alumina particles and seriously react with each other, and leading to the increasing in the rate of dynamic recovery. At last, the interaction between work hardening and working soft-
Table 3 The hardness value of Cu–0.23vol%Al2 O3 alloy compressed under different conditions. Compression conditions
Cu–0.23vol%Al2 O3 alloy
ε˙
0.001 0.01 0.1 1 20
The hardness value (kgf/mm2 ) Original
350 ◦ C
500 ◦ C
650 ◦ C
800 ◦ C
950 ◦ C
120.6
114.9 126.2 127.9 125.2 124.2
100 109.7 116.9 110.7 115.4
96.3 97 107.6 114.8 112.1
72.8 94.6 94.8 103 93
– 83.8 86.5 97.5 91.2
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Fig. 4. Effect of compression conditions on the hardness of Cu–0.23vol%Al2 O3 alloy. (a) Effect of strain rate on the hardness. (b) Effect of temperature on the hardness.
Fig. 5. The original structure of low concentration Cu–0.23vol%Al2 O3 alloy used for thermal simulation. (a) Longitudinal direction and (b) transverse direction.
ening will inevitably leads to the appearance of peak hardness value in a certain strain rate. In addition, as can be seen from Fig. 4(b), with the increasing of compression temperature, the occurrence of dynamic recovery or dynamic recrystallization becomes more and more significant, therefore, a downward trend of hardness for the Cu–0.23vol%Al2 O3 alloy compressed in the same strain rate can be observed.
3.4. Effect of compression conditions on the microstructure In order to explain the change characteristics of the true stress–true strain curves and hardness of the Cu–0.23vol%Al2 O3 alloy compressed under different conditions, it is quite necessary to study the effect of compression conditions on its microstructure. The original microstructure of the Cu–0.23vol%Al2 O3 alloy is
Fig. 6. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 350 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
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Fig. 7. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
Fig. 8. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 650 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
shown in Fig. 5. Highly elongated fibre structure can be observed in the longitudinal direction, and these fibres are arranged to the drawing direction (Fig. 5(a)). Yet, a granular microstructure appears on the transverse cross-section (Fig. 5(b)). The metallographic microstructures of Cu–0.23vol%Al2 O3 alloy compressed under different conditions are shown in Figs. 6–9. From Fig. 6, it can be seen that, the original fibre microstructure is greatly weakened, especially for the case of strain rate 0.001 s−1 yet, dynamic recrystallization grains cannot be observed. When the compression temperature is increased to 500 ◦ C, a large number of fine
dynamic recrystallization grains are formed after compression by three strain rates (as shown in Fig. 7). The most dynamic recrystallization grains can be observed on the overall matrix for the case of strain rate 0.001 s−1 (Fig. 7(a)). With the increasing of strain rate, its deformation becomes more and more non-homogeneous, and length–width ratio of fibre is also increased (Fig. 7(c)). In addition, though fibre structure still can be seen for the case of the higher strain rate (20 s−1 ), yet, most of those fibre structure have been penetrated by the dynamic recrystallization grains thoroughly (Fig. 7(c)). When the compression temperature is increased
Fig. 9. Microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 800 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
K. Shen et al. / Journal of Alloys and Compounds 479 (2009) 401–408
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Fig. 10. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
Fig. 11. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 800 ◦ C. (a) ε˙ = 0.001 s−1 , (b) ε˙ = 0.1 s−1 and (c) ε˙ = 20 s−1 .
to 650 ◦ C, the size of dynamic recrystallization grains is significantly increased (Fig. 8(a)), and the original fibre structure has been divided into the smaller grain cells by dynamic recrystallization grains (as arrow directions in Fig. 8(b) and (c)). When the temperature is further increased to 800 ◦ C, from Fig. 9, we can see that the size of dynamic recrystallization grain is also increased greatly, especially for the case of strain rate 0.001 s−1 . For the different high temperature compression process, both the interaction between dislocation and dispersion particles, and subgrain structure are different for the Cu–0.23vol%Al2 O3 alloy. Therefore, there is need to analyze their TEM microstructure under different conditions. TEM microstructure of low concentration Cu–0.23vol%Al2 O3 alloy compressed at 500 ◦ C with different strain rates is shown in Fig. 10. From it, we can see that, with the increasing of strain rate, the size of subgrain with an equiaxed shape is gradually decreased. When the sample is compressed in the strain rate 0.001 s−1 , subgrain size is the largest (about 2–3 m), and the density of dislocation is much lower, only those dispersion particles distributed in some coarser subgrains are tangled by a few dislocation lines (Fig. 10(a)). With the increasing of strain rate to 0.1 s−1 , a reduced subgrain size can be seen (about 1.5–2 m), yet, the density of dislocation is much higher than that of in strain rate 0.001 s−1 , and most of these dislocation lines mainly distribute around the subgrain boundary. Besides, both the combination of some subgrains, and transformation of small-angle to wide-angle grain boundary can be observed in Fig. 10(b). With further increasing of strain rate (to 20 s−1 ), it can be see that the subgrain size is reduced to only about 0.5–1 m at this moment, and the clear
subgrain boundary can be also observed, yet, the density of dislocation is lower than that of in strain rate 0.1 s−1 , which suggests the higher strain rate is favorable to the formation of subgrain structure. When the temperature is increased to 800 ◦ C, the size of subgrain is much larger than that of compressed at 500 ◦ C in the same strain rate, and the higher strain rate is also favorable to the formation of clear subgrain boundary (Fig. 11), which is the same as that of the lower temperature compression of the Cu–0.23vol%Al2 O3 alloy. 4. Conclusions The properties and microstructure of the Cu–0.23vol%Al2 O3 alloy compressed at high temperatures have been investigated. The results are as follows: (1) The curves of true stress–true strain for the Cu–0.23vol%Al2 O3 alloy compressed at high temperatures can be mainly characterized by an initial sharp increase stage, steady stage, and flow stress softening stage. Both the peak yield stress and strain rate satisfies the hyperbolic sine relationship, indicating its high temperature compression being a thermal activation process. Based on the calculated material parameters, the constitutive equation for peak yield stress, strain rate and temperature was established as follows: ε˙ = [sinh(0.0124836)]4.39909 exp(11.65218 − 99.848 × 103 /RT ). (2) The hardness of the Cu–0.23vol%Al2 O3 alloy compressed at the same temperature is first increased with the increasing of strain
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rate, then after reaching the peak value it begins to decrease with further increasing of strain rate. With the increasing of compression temperature, a downward trend of hardness for this alloy compressed in the same strain rate can be observed. (3) With the increasing of compression temperature, the size and number of dynamic recrystallization grains in the alloy matrix are increased. Yet, when deformed at the same temperature, with the increasing of strain rates, uneven distribution of metallographical microstructure is also strengthened, and the size of subgrain is gradually decreased; yet, the dislocation density is increased at first, and then followed by a decreasing. The higher strain rate is favorable to the formation of fine subgrains with a clear subgrain boundary. Acknowledgement This study was supported by the National High-Tech Research and Development program of China (863 Program).
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