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High creep resistance behavior of the cast Al–Cu alloy modified by nano-scale PrxOy
Materials Science and Engineering A 515 (2009) 10–13
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
High creep resistance behavior of the cast Al–Cu alloy modified by nano-scale Prx Oy W.G. Zhao, J.G. Wang, H.L. Zhao, D.M. Yao, Q.C. Jiang ∗ Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, No. 5988 Renmin Street, Changchun 130025, PR China
a r t i c l e
i n f o
Article history: Received 5 November 2008 Received in revised form 3 February 2009 Accepted 26 March 2009 Keywords: Al–Cu alloys Creep Dislocation movement
a b s t r a c t Creep behavior of the unmodified and modified cast Al–Cu alloys was investigated at temperatures from 393 to 483 K in the tension test. The creep resistance ability of the Al–Cu alloy modified by Prx Oy is almost 3–5 times as high as that of the unmodified Al–Cu alloy, which is attributed to a large number of nanoscale precipitates with high thermal stability in the modified Al–Cu alloy restricting and impeding the dislocation movement during the creep. The induction of a threshold stress in the analysis leads to a stress exponent of 5, which suggests that the creep behavior of both the present alloys is associated with the lattice diffusion-controlled dislocation climb (n = 5). © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Al alloys are becoming increasingly significant as lightweight metal structural materials in many industries, such as aircraft construction, space technology, military field and automobile manufacturing. To magnify the further application of Al alloys at high temperature, thus it is necessary to understand the creep behavior of the Al alloys. Unfortunately, only limited creep investigations reported to date have been conducted on the particle reinforced Al matrix composites and several Al alloys [1–6]. Li and Langdon investigated the creep behavior of an Al-6061 matrix alloy [4]. Recently, Takagi reported the creep characterization of Al–Mg solid-solution alloy [7]. Chauhan studied the creep behavior of near-nanostructured Al 5083 alloy [8]. However, the creep behavior of cast Al alloys around 473 K, especially for the cast Al–Cu alloys has not received much attention. Therefore, the cast Al–Cu alloys in the present study were subjected to the tensile creep investigations at 393–483 K, and five orders of magnitude of creep strain rates were measured. The objective of this study is to probe into the reason for higher creep resistance ability of the present Al–Cu alloy modified by nano-scale Prx Oy than that of the unmodified Al–Cu alloy, based on their creep behavior.
The compositions (measured by an ARL 4460 Metals Analyzer) of the cast Al–Cu alloys were (in wt.%) 6.0 Cu, 0.15 Mn, 0.25 Ti, 0.13 V, 0.13 Zr, 0.001 B and balance Al. The present Al–Cu alloys were modified by nano-scale Prx Oy . Details of the fabrication process for the present Al–Cu alloys are described elsewhere [9]. The heat-treated cast Al–Cu alloys were cut into the tensile dog-bone shaped specimens with a gauge cross-section of 5.0 mm × 2.0 mm and a gauge length of 20.0 mm. Specimen surfaces were polished to a mirror-like finish surface before the creep tests. The creep tests were performed on the tester (CSS-2905, Changchun, China) at 393, 423, 453 and 483 K under the load and stroke control modes with applied stresses from 125 to 250 MPa according to about 0.67–0.96 0.2 ( 0.2 is the tensile yield strength) of the present alloys in the range of investigated temperatures. The temperature of the sample was monitored using a thermocouple tied to the middle of gauge length of the specimen. The specimen temperature was controlled within ±1 K. Microstructures were examined using a transmission electron microscope (TEM) (JEM-2000FX, Tokyo, Japan).
∗ Corresponding author. Fax: +86 431 85094699. E-mail address: [email protected] (Q.C. Jiang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.03.080
3. Results and discussion Fig. 1(a) indicates that the comparison between the creep rate of the unmodified and modified Al–Cu alloys (designated as UA sample and MA sample, respectively) at 453 and 483 K as a function of applied stress. Clearly, the creep rate of the UA sample is almost 3–5 times as high as that of the MA sample. This indicates
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
11
Fig. 2. Variation of the measured minimum creep rate with the applied stress for the MA and UA samples.
Fig. 1. (a) Comparison between creep rates of the unmodified and modified Al–Cu alloys at 453 and 483 K as a function of applied stress; (b) the typical creep curve for the UA and MA samples at an applied stress of 175 MPa at 483 K.
that the MA sample exhibits superior creep resistance over the UA sample in the range of investigated temperatures and creep rates. Besides, in our previous study [9], the MA sample exhibits much higher strength and better ductility (the strength of 520 MPa and ductility of 13.5%) than the UA sample at room temperature. A generalization for above, the MA sample has higher strength, ductility and creep resistance than the UA sample due to Prx Oy addition. Fig. 1(b) shows the typical creep curve for the MA and UA samples obtained at an applied stress of 175 MPa at 483 K. The creep curve shows the occurrence of a short primary stage of creep followed by the secondary stage. In the secondary stage the creep rate is a constant—reaching a steady state, and the creep rate accelerates in the final stage, leading to fracture. Furthermore, the total failure strains for the UA sample are much higher than those for the MA sample under the similar testing conditions. Fig. 2 reveals the variation of the measured minimum creep rate (ε˙ m ), with the applied stress (), for the MA and UA samples at 393, 423, 453 and 483 K. It is noted that the creep data measured for the MA sample cover five orders of magnitude of creep strain rates (2.26 × 10−10 –2.33 × 10−5 s−1 ). Apparently, the data at each temperature in Fig. 2 fit into a straight line. The stress exponent (n) is related to the creep rate and applied stress from the following power-law equation [10,11]: ε˙ m =
AGbD0 kT
n G
Q app
exp −
RT
mann’s constant, is the applied stress, n is the apparent stress exponent, Qapp is the apparent activation energy for creep, R is the gas constant, and T is the absolute temperature. The MA and UA samples exhibit the apparent stress exponents (n) of 18.5, 15.2, 15.0, 12.4 and 13.1, 15.2, 17.2, 12.9 at 393, 423, 453 and 483 K, respectively. This indicates that the apparent stress exponent of the MA sample decreases, while that of the UA sample first increases and then decreases with increasing temperature. The reason for this needs further understanding. Fig. 3 displays the effect of temperature on the steady-state creep rate for both the MA and UA samples. The Qapp (Qapp = ˙ −R(∂ ln ε/∂(1/T ))) was determined to be 173.5 and 150.2 kJ/mol for the MA sample and UA sample, respectively, at a constant stress of 175 MPa. The apparent activation energy for the MA and UA samples estimated in the present investigation was higher than the value either for self-diffusion in Al (∼143.4 kJ/mol [12]) or for diffusion of Cu in Al (∼137 kJ/mol [13]). As discussed above, the apparent stress exponents and creep activation energies of the MA and UA samples are 12.4–18.5, 173.5 kJ/mol and 12.9–17.2, 150.2 kJ/mol, respectively. The values of the apparent stress exponent are much higher than those of the cast pure Al or the solid-solution Al alloys (n = 3–5). The results in the present study indicate that the creep behavior of the MA and UA alloys is analogous to that of the dispersion strengthened (DS) alloys [14].
(1)
where ε˙ m is the minimum creep rate, A is a structure dependent constant, G is the temperature-dependent shear modulus, b is the Burgers vector, D0 is the frequency factor for diffusion, k is the Boltz-
Fig. 3. Effect of temperature on the steady-state creep rate of both the UA and MA samples.
12
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
1/n
Fig. 4. Plots of ε˙ m vs. for n = (a) 2, (b) 3 and (c) 5 on linear scales for the MA sample.
The high values of apparent stress exponent and apparent activation energy in the DS alloys have been explained by means of introducing a threshold stress [15]. When creep occurs in the presence of a threshold stress, the minimum or steady-state creep rate may be expressed in the form [4]: ε˙ m =
AGbD0 kT
− n 0
G
Q
exp −
RT
(2)
where 0 is the threshold stress, n is the true stress exponent, Q is the true activation energy for creep, the rest symbols in Eq. (2) have the same meanings as indicated in Eq. (1). Assuming that the creep data of the present alloys satisfy Eq. (2) and 0 is independent of the applied stress, if the data points in 1/n the plot of ε˙ m vs. fit a straight line, the value of 0 at each temperature can be determined by extrapolating the linear regression line to a zero creep rate. A set of true stress exponent, n = 2, 3, and 5
Fig. 5. TEM micrograph of the UA (a), MA (c) samples before the creep test, and the UA (b), MA (d) samples after the creep test at 483 K and a stress of 150 MPa.
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
were used to fit the experimental data points. These values of n correspond to three well-documented creep cases for metal and alloys: n = 2 for creep controlled by the grain boundary sliding [16,17], n = 3 for creep controlled by the dislocation viscous glide process [17,18], and n = 5 for creep controlled by the high-temperature dislocation climb [17,18]. 1/n Fig. 4 shows the plots of ε˙ m vs. for n = 2, 3 and 5 on linear scales for the MA sample. The results reveal that the stress exponent 1/n of 5 yields the best linear fit between ε˙ m vs. ; for n = 2 and 3, the results exhibit a clear curvature with increasing applied stress. Similar result is obtained for the UA sample. This indicates that the creep behavior of the MA and UA samples is not controlled by grain boundary sliding (n = 2) or dislocation viscous glide (n = 3) but controlled by high-temperature dislocation climb. Fig. 5 shows the TEM micrograph of the UA and MA samples before and after the creep test at 483 K and a stress of 150 MPa. There were only a few needle-shaped precipitates with the width of 10–15 nm and the length of 120–140 nm in the UA sample in Fig. 5(a); while, there were a large number of needle-shaped precipitates with the width of 8–10 nm and the length of 100–120 nm in the MA sample in Fig. 5(c). At the same time, the distribution of the nano-scale precipitates in the MA sample was regular and homogeneous. These results are in very good agreement with our previous study [9]. This deduces that nano-scale Prx Oy addition facilitates the formation of precipitates during the ageing heattreatment, which is responsible for the high creep resistance ability of the MA sample during creep test. The present creep experiments were conducted at temperatures close to or higher than the peak ageing temperature of the present alloys, thereby introducing the possibility of partial dissolution and coarsening of some precipitates (such as precipitates). As shown in Fig. 5(b), the dissolution and growth of precipitates in the UA sample, together with some precipitates coarsening (the width from 10–15 to 20–35 nm and the length from 120–140 to 150–180 nm) is obvious. While, in the MA sample (indicated in Fig. 5(d), the growth of precipitates as well as precipitates coarsening is not evident, although some precipitates grow and coarsen (the width from 8–10 to 15–25 nm and the length from 100–120 to 120–140 nm). Besides, the amount of precipitates in the MA sample is more than that in the UA sample during the creep test. This result indicates that Prx Oy addition improves the thermal stability of precipitates in the MA sample during the creep test at high temperature, which is responsible for higher creep resistance of the MA sample. Further work is continuing in order to completely understand the mechanism for Prx Oy addition enhancing the thermal stability of the precipitates in the MA sample. It is well established that the high-temperature dislocation climb model predicts a stress exponent of 5 for lattice-diffusion controlled creep [18]. However, if the creep data of the present MA
13
and UA samples are described by the high-temperature dislocation climb model together with a threshold stress, the values of true activation energy of 130.7 and 120.5 kJ/mol for the MA and UA samples are obtained, respectively. Obviously, the value of Q for the UA sample is lower than that for the MA sample. As mentioned above, the thermal stability of precipitates in the MA sample is higher than that in the UA sample according to the analyses of TEM microstructures in Fig. 5. Thus, a large number of nano-scale precipitates with high thermal stability in the MA sample restrict and impede the dislocations movement during the creep test at high temperatures resulting in the enhanced creep resistance ability of the MA sample. 4. Conclusion The creep resistance ability of the MA sample is almost 3–5 times higher than that of the UA sample, which is attributed to a large number of nano-scale precipitates with high thermal stability in the MA sample restricting and impeding the dislocation movement during creep. It is demonstrated, by incorporating a threshold stress in the analysis, that the true stress exponent is equal to 5, which suggests that the creep behavior of the UA and MA samples is associated with the lattice diffusion-controlled dislocation climb. Acknowledgements This work is supported by The National Natural Science Foundation of China (No. 50771050) and The Project 985-Automotive Engineering of Jilin University. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
T.G. Nieh, Metall. Trans. A 15 (1984) 139–146. K.T. Park, E.J. Lavernia, F.A. Mohamed, Acta Metall. Mater. 38 (1990) 2149–2159. A.B. Pandey, R.S. Mishra, Y.R. Mahajan, Acta Metall. Mater. 40 (1992) 2045–2052. Y. Li, T.G. Langdon, Acta Mater. 45 (1997) 4797–4806. S.C. Tjong, Z.Y. Ma, Mater. Sci. Technol. 59 (1999) 1117–1125. A.B. Pandey, R.S. Mishra, Y.R. Mahajan, Scripta Metall. Mater. 24 (1990) 1565–1570. H. Takagi, M. Dao, M. Fujiwara, M. Otsuka, Mater. Trans. 47 (2006) 2006–2014. M. Chauhan, I. Roy, F.A. Mohamed, Mater. Sci. Eng. A 410 (2005) 24–27. W.G. Zhao, H.Y. Wang, J.G. Wang, Y. Li, Q.C. Jiang, J. Mater. Res. 23 (2008) 1076–1081. Z.Y. Ma, S.C. Tjong, X.M. Meng, J. Mater. Res. 17 (2002) 2307–2313. J. Weertman, J. Appl. Phys. 26 (1955) 1213–1217. F.A. Mohamed, T.G. Langdon, Metall. Trans. 5 (1974) 2339–2345. I.S. Grigoriev, E.Z. Meilikhov, Handbook of Physical Quantities, CRC Press, New York, 1997. L. Kloc, S. Spigarelli, E. Evangelista, T.G. Langdon, Acta Mater. 45 (1997) 529–540. S. Purushothaman, J.K. Tien, Acta Metall. 26 (1978) 519. R.C. Gifkins, Metall. Mater. Trans. A 7 (1976) 1225–1232. P. Zhang, Scripta Mater. 52 (2005) 277–282. F.A. Mohamed, T.G. Langdon, Acta Metall. 22 (1974) 779–788.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
High creep resistance behavior of the cast Al–Cu alloy modified by nano-scale Prx Oy W.G. Zhao, J.G. Wang, H.L. Zhao, D.M. Yao, Q.C. Jiang ∗ Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, No. 5988 Renmin Street, Changchun 130025, PR China
a r t i c l e
i n f o
Article history: Received 5 November 2008 Received in revised form 3 February 2009 Accepted 26 March 2009 Keywords: Al–Cu alloys Creep Dislocation movement
a b s t r a c t Creep behavior of the unmodified and modified cast Al–Cu alloys was investigated at temperatures from 393 to 483 K in the tension test. The creep resistance ability of the Al–Cu alloy modified by Prx Oy is almost 3–5 times as high as that of the unmodified Al–Cu alloy, which is attributed to a large number of nanoscale precipitates with high thermal stability in the modified Al–Cu alloy restricting and impeding the dislocation movement during the creep. The induction of a threshold stress in the analysis leads to a stress exponent of 5, which suggests that the creep behavior of both the present alloys is associated with the lattice diffusion-controlled dislocation climb (n = 5). © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Al alloys are becoming increasingly significant as lightweight metal structural materials in many industries, such as aircraft construction, space technology, military field and automobile manufacturing. To magnify the further application of Al alloys at high temperature, thus it is necessary to understand the creep behavior of the Al alloys. Unfortunately, only limited creep investigations reported to date have been conducted on the particle reinforced Al matrix composites and several Al alloys [1–6]. Li and Langdon investigated the creep behavior of an Al-6061 matrix alloy [4]. Recently, Takagi reported the creep characterization of Al–Mg solid-solution alloy [7]. Chauhan studied the creep behavior of near-nanostructured Al 5083 alloy [8]. However, the creep behavior of cast Al alloys around 473 K, especially for the cast Al–Cu alloys has not received much attention. Therefore, the cast Al–Cu alloys in the present study were subjected to the tensile creep investigations at 393–483 K, and five orders of magnitude of creep strain rates were measured. The objective of this study is to probe into the reason for higher creep resistance ability of the present Al–Cu alloy modified by nano-scale Prx Oy than that of the unmodified Al–Cu alloy, based on their creep behavior.
The compositions (measured by an ARL 4460 Metals Analyzer) of the cast Al–Cu alloys were (in wt.%) 6.0 Cu, 0.15 Mn, 0.25 Ti, 0.13 V, 0.13 Zr, 0.001 B and balance Al. The present Al–Cu alloys were modified by nano-scale Prx Oy . Details of the fabrication process for the present Al–Cu alloys are described elsewhere [9]. The heat-treated cast Al–Cu alloys were cut into the tensile dog-bone shaped specimens with a gauge cross-section of 5.0 mm × 2.0 mm and a gauge length of 20.0 mm. Specimen surfaces were polished to a mirror-like finish surface before the creep tests. The creep tests were performed on the tester (CSS-2905, Changchun, China) at 393, 423, 453 and 483 K under the load and stroke control modes with applied stresses from 125 to 250 MPa according to about 0.67–0.96 0.2 ( 0.2 is the tensile yield strength) of the present alloys in the range of investigated temperatures. The temperature of the sample was monitored using a thermocouple tied to the middle of gauge length of the specimen. The specimen temperature was controlled within ±1 K. Microstructures were examined using a transmission electron microscope (TEM) (JEM-2000FX, Tokyo, Japan).
∗ Corresponding author. Fax: +86 431 85094699. E-mail address: [email protected] (Q.C. Jiang). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.03.080
3. Results and discussion Fig. 1(a) indicates that the comparison between the creep rate of the unmodified and modified Al–Cu alloys (designated as UA sample and MA sample, respectively) at 453 and 483 K as a function of applied stress. Clearly, the creep rate of the UA sample is almost 3–5 times as high as that of the MA sample. This indicates
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
11
Fig. 2. Variation of the measured minimum creep rate with the applied stress for the MA and UA samples.
Fig. 1. (a) Comparison between creep rates of the unmodified and modified Al–Cu alloys at 453 and 483 K as a function of applied stress; (b) the typical creep curve for the UA and MA samples at an applied stress of 175 MPa at 483 K.
that the MA sample exhibits superior creep resistance over the UA sample in the range of investigated temperatures and creep rates. Besides, in our previous study [9], the MA sample exhibits much higher strength and better ductility (the strength of 520 MPa and ductility of 13.5%) than the UA sample at room temperature. A generalization for above, the MA sample has higher strength, ductility and creep resistance than the UA sample due to Prx Oy addition. Fig. 1(b) shows the typical creep curve for the MA and UA samples obtained at an applied stress of 175 MPa at 483 K. The creep curve shows the occurrence of a short primary stage of creep followed by the secondary stage. In the secondary stage the creep rate is a constant—reaching a steady state, and the creep rate accelerates in the final stage, leading to fracture. Furthermore, the total failure strains for the UA sample are much higher than those for the MA sample under the similar testing conditions. Fig. 2 reveals the variation of the measured minimum creep rate (ε˙ m ), with the applied stress (), for the MA and UA samples at 393, 423, 453 and 483 K. It is noted that the creep data measured for the MA sample cover five orders of magnitude of creep strain rates (2.26 × 10−10 –2.33 × 10−5 s−1 ). Apparently, the data at each temperature in Fig. 2 fit into a straight line. The stress exponent (n) is related to the creep rate and applied stress from the following power-law equation [10,11]: ε˙ m =
AGbD0 kT
n G
Q app
exp −
RT
mann’s constant, is the applied stress, n is the apparent stress exponent, Qapp is the apparent activation energy for creep, R is the gas constant, and T is the absolute temperature. The MA and UA samples exhibit the apparent stress exponents (n) of 18.5, 15.2, 15.0, 12.4 and 13.1, 15.2, 17.2, 12.9 at 393, 423, 453 and 483 K, respectively. This indicates that the apparent stress exponent of the MA sample decreases, while that of the UA sample first increases and then decreases with increasing temperature. The reason for this needs further understanding. Fig. 3 displays the effect of temperature on the steady-state creep rate for both the MA and UA samples. The Qapp (Qapp = ˙ −R(∂ ln ε/∂(1/T ))) was determined to be 173.5 and 150.2 kJ/mol for the MA sample and UA sample, respectively, at a constant stress of 175 MPa. The apparent activation energy for the MA and UA samples estimated in the present investigation was higher than the value either for self-diffusion in Al (∼143.4 kJ/mol [12]) or for diffusion of Cu in Al (∼137 kJ/mol [13]). As discussed above, the apparent stress exponents and creep activation energies of the MA and UA samples are 12.4–18.5, 173.5 kJ/mol and 12.9–17.2, 150.2 kJ/mol, respectively. The values of the apparent stress exponent are much higher than those of the cast pure Al or the solid-solution Al alloys (n = 3–5). The results in the present study indicate that the creep behavior of the MA and UA alloys is analogous to that of the dispersion strengthened (DS) alloys [14].
(1)
where ε˙ m is the minimum creep rate, A is a structure dependent constant, G is the temperature-dependent shear modulus, b is the Burgers vector, D0 is the frequency factor for diffusion, k is the Boltz-
Fig. 3. Effect of temperature on the steady-state creep rate of both the UA and MA samples.
12
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
1/n
Fig. 4. Plots of ε˙ m vs. for n = (a) 2, (b) 3 and (c) 5 on linear scales for the MA sample.
The high values of apparent stress exponent and apparent activation energy in the DS alloys have been explained by means of introducing a threshold stress [15]. When creep occurs in the presence of a threshold stress, the minimum or steady-state creep rate may be expressed in the form [4]: ε˙ m =
AGbD0 kT
− n 0
G
Q
exp −
RT
(2)
where 0 is the threshold stress, n is the true stress exponent, Q is the true activation energy for creep, the rest symbols in Eq. (2) have the same meanings as indicated in Eq. (1). Assuming that the creep data of the present alloys satisfy Eq. (2) and 0 is independent of the applied stress, if the data points in 1/n the plot of ε˙ m vs. fit a straight line, the value of 0 at each temperature can be determined by extrapolating the linear regression line to a zero creep rate. A set of true stress exponent, n = 2, 3, and 5
Fig. 5. TEM micrograph of the UA (a), MA (c) samples before the creep test, and the UA (b), MA (d) samples after the creep test at 483 K and a stress of 150 MPa.
W.G. Zhao et al. / Materials Science and Engineering A 515 (2009) 10–13
were used to fit the experimental data points. These values of n correspond to three well-documented creep cases for metal and alloys: n = 2 for creep controlled by the grain boundary sliding [16,17], n = 3 for creep controlled by the dislocation viscous glide process [17,18], and n = 5 for creep controlled by the high-temperature dislocation climb [17,18]. 1/n Fig. 4 shows the plots of ε˙ m vs. for n = 2, 3 and 5 on linear scales for the MA sample. The results reveal that the stress exponent 1/n of 5 yields the best linear fit between ε˙ m vs. ; for n = 2 and 3, the results exhibit a clear curvature with increasing applied stress. Similar result is obtained for the UA sample. This indicates that the creep behavior of the MA and UA samples is not controlled by grain boundary sliding (n = 2) or dislocation viscous glide (n = 3) but controlled by high-temperature dislocation climb. Fig. 5 shows the TEM micrograph of the UA and MA samples before and after the creep test at 483 K and a stress of 150 MPa. There were only a few needle-shaped precipitates with the width of 10–15 nm and the length of 120–140 nm in the UA sample in Fig. 5(a); while, there were a large number of needle-shaped precipitates with the width of 8–10 nm and the length of 100–120 nm in the MA sample in Fig. 5(c). At the same time, the distribution of the nano-scale precipitates in the MA sample was regular and homogeneous. These results are in very good agreement with our previous study [9]. This deduces that nano-scale Prx Oy addition facilitates the formation of precipitates during the ageing heattreatment, which is responsible for the high creep resistance ability of the MA sample during creep test. The present creep experiments were conducted at temperatures close to or higher than the peak ageing temperature of the present alloys, thereby introducing the possibility of partial dissolution and coarsening of some precipitates (such as precipitates). As shown in Fig. 5(b), the dissolution and growth of precipitates in the UA sample, together with some precipitates coarsening (the width from 10–15 to 20–35 nm and the length from 120–140 to 150–180 nm) is obvious. While, in the MA sample (indicated in Fig. 5(d), the growth of precipitates as well as precipitates coarsening is not evident, although some precipitates grow and coarsen (the width from 8–10 to 15–25 nm and the length from 100–120 to 120–140 nm). Besides, the amount of precipitates in the MA sample is more than that in the UA sample during the creep test. This result indicates that Prx Oy addition improves the thermal stability of precipitates in the MA sample during the creep test at high temperature, which is responsible for higher creep resistance of the MA sample. Further work is continuing in order to completely understand the mechanism for Prx Oy addition enhancing the thermal stability of the precipitates in the MA sample. It is well established that the high-temperature dislocation climb model predicts a stress exponent of 5 for lattice-diffusion controlled creep [18]. However, if the creep data of the present MA
13
and UA samples are described by the high-temperature dislocation climb model together with a threshold stress, the values of true activation energy of 130.7 and 120.5 kJ/mol for the MA and UA samples are obtained, respectively. Obviously, the value of Q for the UA sample is lower than that for the MA sample. As mentioned above, the thermal stability of precipitates in the MA sample is higher than that in the UA sample according to the analyses of TEM microstructures in Fig. 5. Thus, a large number of nano-scale precipitates with high thermal stability in the MA sample restrict and impede the dislocations movement during the creep test at high temperatures resulting in the enhanced creep resistance ability of the MA sample. 4. Conclusion The creep resistance ability of the MA sample is almost 3–5 times higher than that of the UA sample, which is attributed to a large number of nano-scale precipitates with high thermal stability in the MA sample restricting and impeding the dislocation movement during creep. It is demonstrated, by incorporating a threshold stress in the analysis, that the true stress exponent is equal to 5, which suggests that the creep behavior of the UA and MA samples is associated with the lattice diffusion-controlled dislocation climb. Acknowledgements This work is supported by The National Natural Science Foundation of China (No. 50771050) and The Project 985-Automotive Engineering of Jilin University. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
T.G. Nieh, Metall. Trans. A 15 (1984) 139–146. K.T. Park, E.J. Lavernia, F.A. Mohamed, Acta Metall. Mater. 38 (1990) 2149–2159. A.B. Pandey, R.S. Mishra, Y.R. Mahajan, Acta Metall. Mater. 40 (1992) 2045–2052. Y. Li, T.G. Langdon, Acta Mater. 45 (1997) 4797–4806. S.C. Tjong, Z.Y. Ma, Mater. Sci. Technol. 59 (1999) 1117–1125. A.B. Pandey, R.S. Mishra, Y.R. Mahajan, Scripta Metall. Mater. 24 (1990) 1565–1570. H. Takagi, M. Dao, M. Fujiwara, M. Otsuka, Mater. Trans. 47 (2006) 2006–2014. M. Chauhan, I. Roy, F.A. Mohamed, Mater. Sci. Eng. A 410 (2005) 24–27. W.G. Zhao, H.Y. Wang, J.G. Wang, Y. Li, Q.C. Jiang, J. Mater. Res. 23 (2008) 1076–1081. Z.Y. Ma, S.C. Tjong, X.M. Meng, J. Mater. Res. 17 (2002) 2307–2313. J. Weertman, J. Appl. Phys. 26 (1955) 1213–1217. F.A. Mohamed, T.G. Langdon, Metall. Trans. 5 (1974) 2339–2345. I.S. Grigoriev, E.Z. Meilikhov, Handbook of Physical Quantities, CRC Press, New York, 1997. L. Kloc, S. Spigarelli, E. Evangelista, T.G. Langdon, Acta Mater. 45 (1997) 529–540. S. Purushothaman, J.K. Tien, Acta Metall. 26 (1978) 519. R.C. Gifkins, Metall. Mater. Trans. A 7 (1976) 1225–1232. P. Zhang, Scripta Mater. 52 (2005) 277–282. F.A. Mohamed, T.G. Langdon, Acta Metall. 22 (1974) 779–788.