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Effects of Gd and Y additions on microstructure and properties of Al–Zn–Mg–Cu–Zr alloys
Materials Science and Engineering A 552 (2012) 230–235
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effects of Gd and Y additions on microstructure and properties of Al–Zn–Mg–Cu–Zr alloys XingGuo Zhang ∗ , FeiQiang Mei, HuanYue Zhang, ShaoHua Wang, CanFeng Fang, Hai Hao Foundry Center, Dalian University of Technology, Dalian 116024, China
a r t i c l e
i n f o
Article history: Received 29 March 2012 Received in revised form 11 May 2012 Accepted 13 May 2012 Available online 19 May 2012 Keywords: Gd Y Microstructure Mechanical properties
a b s t r a c t Three kinds of alloys Al–Zn–Mg–Cu–Zr, A1–Zn–Mg–Cu–Zr–Gd and A1–Zn–Mg–Cu–Zr–Gd–Y were prepared by cast metallurgy. The effects of Gd and Y additions on the microstructure and properties of Al–Zr–Mg–Cu–Zr alloys were investigated by optical microscopy (OM), tensile test, scanning electron microscopy (SEM), electronic probe microanalysis (EPMA) and transmission electronic microscopy (TEM). The results show that the combined additions of Gd and Y to the A1–Zn–Mg–Cu–Zr alloys effectively retard the mergence and growth of the sub-grain during the recovery process, significantly inhibit the recrystallization of Al matrix. b , 0.2 and ı are improved with increasing Gd and Y content, and the improving amplitudes reach 11.49%, 14.61%, 63.11% respectively and have better relative intracrystalline solubility. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Super-high strength aluminum alloy, which was set against the background of aerospace materials in 1960s, is a kind of high-strength aluminum alloy materials [1–3]. High strength, low density and good mechanical properties make Al–Zn–Mg–Cu alloys attractive for applications in the military and aerospace industries. But with the rapid development of composite material and titanium alloy, ultra-high strength aluminum alloy is facing unprecedented challenges. Improving the general properties of ultra-high strength aluminum alloy has been more and more focusing on microalloying. It was well accepted that trace additions of rare earth elements are of great importance to improve the microstructure and mechanical properties of aluminum alloys [4,5]. Zr, which can improve the strength, fracture toughness and stress corrosion resistance performance of ultra-high strength aluminum alloys, now has become the essential element for Al–Zn–Mg–Cu alloy [6,7]. Sc has been studied mostly as beneficial micro alloying element due to the presence of elastically hard, coherent and nano-sized Al3 Sc particles, which strongly refine as-cast grain size, inhibit the dynamic recrystallization and dislocation movement [8]. Furthermore, combined additions of Sc and Zr are shown to be more effective in refining as-cast grain size and improving mechanical properties because of the newly formed Al3 (Sc, Zr) particles in the Al–Zn–Mg–Cu alloys [9]. Y can effectively refine the branch crystal
∗ Corresponding author. Tel.: +86 0411 84706183; fax: +86 0411 84706183. E-mail address: [email protected] (X. Zhang). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.05.035
microstructure of as-cast alloy, inhibit the formation of bulky eutectic microstructure, and improve the elongation [10,11]. The micro alloying researches of rare earth among the aluminium alloy have developed from single microalloying to composite microalloying. But due to the complexity of the aluminium alloy systems as well as the limit and difficulty of characterizing the effect of rare earth. What is more, there are no complete theories available to guide the design of this type due to the lack of research on it. Setting up the relevant model of rare-earth micro alloying as soon as possible provides a new way to predict and improve the final properties [12]. Under this background, this paper will discuss the effect of the composite microalloying of the rare earth Gd and Y in Al–Zn–Mg–Cu–Zr super-high aluminium alloy, and provide reference to rare-earth micro alloying through cast ingot metallurgy process.
2. Experimental 2.1. Materials and processing The alloys investigated belong to the ultra-high strength Al–Zn–Mg–Cu–Zr alloys and include three kinds of alloys: Al–Zn–Mg–Cu–Zr, Al–Zn–Mg–Cu–Zr alloy with Gd additions (Al–Zn–Mg–Cu–Zr–0.25Gd), Al–Zn–Mg–Cu–Zr alloy with Gd and Y additions (Al–Zn–Mg–Cu–Zr–0.1Gd–0.1Y). The chemical composition of the alloys, analyzed by XRF-1800 type of X-ray fluorescence spectrum analyzer, is given in Table 1. High purity Al (99.99%), Mg (99.9%) and Gd (99.9%), and Al–50%Cu, Al–4%Zr, Mg–40%Y master alloys were taken as the starting raw material. The alloys were
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Table 1 Chemical composition of Al–Zn–Mg–Cu–Zr–Gd–Y alloys (mass fraction, %). Alloy
Zn
Mg
Cu
Zr
Gd
Y
1# 2# 3#
7.36 7.27 7.18
1.68 1.47 1.55
1.66 1.71 1.53
0.14 0.12 0.16
0 0.25 0.13
0 0 0.07
melted in an electrical-resistance furnace and cast into an iron mould.
2.2. Heat treatments The as-cast ingots were homogenized at 400 ◦ C for 4 h plus 460 ◦ C for 24 h, followed by air cooling to the room temperature. Then the ingots were converted into plates, with an rolling ratio of 10:1. Subsequently, the samples were solid solution treated at 470 ◦ C for 2 h and quenched in ambient temperature water immediately (no stretching was applied). Following quenching, the plates were aged at 120 ◦ C for 24 h for T6 tempered.
Fe ≤ 0.05 ≤ 0.05 ≤ 0.05
Si ≤ 0.05 ≤ 0.05 ≤ 0.05
Al Bal. Bal. Bal.
2.4. Microstructural characterizations The microstructure was observed by MEFS multifunction optical microscopy. The fractures were observed by scanning electro microscope (SEM). The thin foils for transmission electron microscopy (TEM) were prepared by twin jet-polishing in electrolyte solution of HNO3 and methyl alcohol (1:3). TEM examinations were performed through a JEM-2010 microscope operated at 200 kV. X-ray diffraction analysis (XRD) was used to characterize the phases of samples on a D/max-2500 diffractometer with Cu K␣ radiation. 3. Results and discussion
2.3. Mechanical properties and relative intracrystalline solubility
3.1. Microstructure of the alloys
The tensile properties were tested at room temperature by an electro universal test machine. The tests were performed at a nominal tension rate of 2 mm/min according to the IGC test standard of GB/T 228-2002. The hardness that can be used to monitor the aging processing was measured by light loaded durometer. The magnitude of load was 1000 kg with a 30 s load time. The relative intracrystalline solubility of main alloy elements in the grain was determined by electron probe microanalysis (EPMA) with EPMA-1600.
The microstructures of the studied alloys are shown in Fig. 1. Coarsening as-cast grains (the average size is 62 m), dendritic structures and serious segregation can be obviously observed in alloy 1# . Adding rare-earth could greatly refine the as-cast grain and abated the dendritic structures. The most effective refining appears in alloy 2# , with the average size 32 m, and most dendritic structures vanish (as shown in Fig. 1(2# )). Compared to alloy 2# , simultaneously adding of Gd and Y, the microstructure appears a certain coarsening, with the average size 40 m, but is still basically equiaxial crystal.
Fig. 1. Optical micrographs of the as-cast microstructure.
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Fig. 2. Optical microstructures of two studied alloys with solution treatment.
With the adding of Gd and Y, crystalline grains of ␣(Al) of the alloys in the as-cast structure are refined, that is related to the characteristic of Gd and Y in the process of solidification of the alloys. The chemical properties of Gd and Y elements are rather active and the electronegativities of Gd and Y are 1.2 and 1.22, respectively, smaller than Al(1.5). They can fill the vacancies on the surfaces when they dissolve in the alloys which can reduce the surface tension between two phases. Consequently, nucleation speed can increase. Meanwhile, active surface membranes form between grains and the liquid alloy which block the growing of the grains and make the grains refined. Grain refinement can be obtained that compared with Al, the atomic radius of Gd and Y are larger. The results show the substitutional solid solution is formed when Gd and Y additions are dissolved into Al [13,14]. When Y is added into Al–Zn–Mg–Cu–Zr–Gd alloy, the effect of grain refinement is weaker than that adding Gd element only. It is mainly due to the interaction among Zr, Gd and Y elements that weaken the effect of grain refinement [11]. Typical optical microstructures of T6-tempered Al–Zn–Mg–Cu–Zr–Gd–Y alloys, which had been rolled, solution strengthened, T6 ageing strengthened and corroded by Kelvin solution, are shown in Fig. 2. We could conclude from the pictures that Al–Zn–Mg–Cu–Zr alloy is recrystallized. Some of the subgrains grow up into large recrystallized grains and some of the subgrains boundary disappear
(Fig. 2(1# )). Al–Zn–Mg–Cu–Zr–0.25Gd alloy is partially recrystallized (Fig. 2(2# )). Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy has a fibrous unrecrystallized microstructure with fine subgrain microstructures (Fig. 2(3# )). It is clear that the recrystallization resistance of T6-tempered Al–Zn–Mg–Cu–Zr alloy can be further enhanced by adding Gd and Y in combination. TEM micrographs and XRD of T6-tempered The Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy are presented in Fig. 3. The 10–20 nm spherical dispersoids of second phase have precipitated from matrix after 120 ◦ C, 24 h aging treatment (as shown in Fig. 3(a) and (b)). The chemical composition of the dispersoids was determined by EDXS. The spherical dispersoids contain Gd, Y, Zn, Cu and Al. Based on the XRD (as shown in Fig. 3(d)) and EDXS, it is suggested that Y was dissolved in Al3 CuGd. The Al3 CuGd particles which can strongly pin dislocations and grain boundary (as shown in Fig. 3(c)) appear in the alloys. Thus these particles block the movement of dislocations as well as subboundary and have strong stabilization effect of the substructure of the deformation organization. 3.2. Relative intracrystalline solubility The EPMA images and distribution of major elements of alloys in 1# , 2# and 3# are shown in Fig. 4. The EPMA image of alloy 1# shows that major elements (Zn, Mg and Cu) in Al–Zn–Mg–Cu–Zr
Fig. 3. The TEM micrographs and XRD of T6-tempered Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy.
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Fig. 4. Mapping analyses of main alloy elements in Al–Zn–Mg–Cu–Zr–Gd–Y alloy billets.
enrich on the boundary of the grain (as shown in Fig. 4(1# )). These results indicate that the concentration of elements on the boundary is higher than that in the grain, especially Cu. This is because during the solidification process of alloys, Cu is vulnerable to segregate in the grain boundary in the form of non-equilibrium eutectic compound or the insoluble second phase particles. In the specimen with no additional rare-earth elements, Zn and Mg are rarely dissolved in the grain, but most of them are eutectic segregation on grain boundary. In alloy 2# and 3# , although there is a large concentration gradient of Cu, the segregation of alloying element Zn and Mg on the boundary reduces which increases the grain relative solid-solubility and reduces the micro-segregation of the alloys. The segregation degree of solute elements of the ingot casting is suppressed greatly. In order to point out the grain relative solid-solubility of major elements and offer reference for the following heat treatment technology, this paper adopt the relative intracrystalline solubility of elements (the relative intracrystalline solubility in the grain of element A = content of element A in grain/content of alloying element A) to represent the effect of different contents of rare earth content on the solid-solubility [15]. The results are shown in Table 2. The micro-segregation of the Al–Zn–Mg–Cu–Zr alloy could be significantly reduced after adding Gd or both with Gd and Y. Meanwhile, the relative intracrystalline solubility of Zn, Mg and Cu elements could be improved, that is: from 54.82%, 49.63%, 26.38% to 69.78%,
78.75%, 42.19% after adding 0.25% Gd, and to 69.57%, 79.15%, 49.73% after adding 0.2% combined Gd and Y respectively. According to modern statistical physics, equilibrium partition coefficient can be expressed in diffusion activation energy [16]: k0 = exp
(Q − Q ) L S KT
where QS is diffusion activation energy which the solute atoms needed to across the interface barrier layer into the melt; QL is diffusion activation energy which the solute atoms needed to across the interface barrier layer into the solid solution; K is Boltzmann constant; T is phase transition temperature. Distribution coefficient of solute essentially reflects the difference that solute atoms through the solid-liquid boundary layer to the sides on micro-interface boundary. The equilibrium partition coefficient of Gd and Y are less than 1 and the solid solubility in the Table 2 Relative intracrystalline solubility in the grain of main alloy elements. Alloy
1# 2# 3#
Relative intracrytalline solubility Zn
Mg
Cu
54.82 69.78 69.57
49.63 78.75 79.15
26.38 42.19 49.73
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Fig. 5. Ageing–hardening curves of artificial ageing of alloys.
aluminum is very small, so most of them will enrich in the liquid boundary, which reduces the chance of Zn, Mg and Cu moving into the solid solution and also increases the solid solubility of Zn, Mg and Cu in grain, reduces the microsegregation of alloy. 3.3. Age-hardening curves of the alloys The artificial age-hardening curves of alloys are shown in Fig. 5. It indicates that the aging strengthening properties of these studied alloys are obviously improved with prolonging aging time. The improvements of peak-aging hardness are all above 50% and the artificial age-hardening curves of the alloy 2# and alloy 3# stay consistent approximately. In the process of artificial age-hardening, the peak-aging of alloy 1# locates at 11–24 h while alloy 2# and alloy 3# appear after 7 h and last up to 32 h. Compared to the alloy 1# , the
peak-aging hardness of the alloy 2# and the alloy 3# improve about 20 HBS. The results show that the addition of rare earth extends the peak-aging time, makes the peak-aging arriving earlier and improves the peak-aging hardness of artificial age-hardening at 120 ◦ C. The mainly strengthened phase in Al–Zn–Mg(–Cu) alloy is (MgZn2 ) and , and the general precipitation sequence is as follows: a sosoloid → G.P. zones → (MgZn2 ) → phase. Two kinds of effects of Gd and Y can be observed in Al–Zn–Mg–Cu–Zr alloys. One is minor rare earth element addition can improve the concentration of Zn, Mg and Cu into a sosoloid, as suggested before, therefore the driving force of phase transition is enhanced from G.P. zones to (MgZn2 ) phase [17]. The other is Al3 CuGd particles can come into being nucleation core during aging when these particles do not melt back after solution treatment and distribute dispersively in the alloy matrix. Thus minor rare earth element addition can make the peak aging arriving earlier for Al–Zn–Mg–Cu–Zr alloys. Gd and Y are elements that display slightly lower binding energy with vacancies than are exhibited by Mg and Zn [18,19], which will automatically preferentially attach to vacancies and sequester them. Such a preferential binding of vacancies with Gd and Y or both results in a decrease in the number of vacancies available for the movement of Mg and Zn atoms and hence is exhibited in delayed growth of the ’ (MgZn2 ) phase. Thus the observed delay in reaching the peak after adding Gd only or both with Gd and Y is directly related to the binding energy effect. 3.4. Tensile properties of the alloys under T6 treatment The tensile properties of the alloys under T6 treatment are shown in Table 3. It indicates that compared with non-RE addition alloy, adding Gd only or both with Gd and Y can significantly improve the strength and elongation of T6 state. When the alloy containing 0.25% Gd, the ultimate tensile strength and yield strength values are 624.54 MPa, 595.00 MPa at T6 and the
Fig. 6. SEM micrographs of tensile fracture surface of alloys.
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Fig. 7. Morphology and composition of second-phase in the dimple.
Table 3 Mechanical properties of alloys. Alloy #
1 2# 3#
b (MPa)
0.2 (MPa)
ı (%)
573.44 624.54 639.30
544.30 595.00 623.80
10.3 13.3 16.8
elongation is 13.3%. The improving amplitudes reach 8.91%, 9.31%, 29.13%, respectively. When adding both compound Gd and Y elements, the properties are optimal: tensile strength is up to 639.3 Mpa, yield strength is 623.8 Mpa, elongation is 16.8%, and the improving amplitudes reach 11.49%, 14.61%, 63.11%, respectively. Fig. 6 shows the SEM fracture morphology of the tensile test specimen under T6 treatment. The main fracture behavior of Al–Zn–Mg–Cu–Zr alloy with no additional rare-earth elements is shear fracture and dimple intercrystalline fracture (as shown in Fig. 6(1# )), whereas the primary fracture mechanism of the alloy with additional rare-earth elements is dimpled transgranular type of failure (as shown in Fig. 6(2# ), Fig. 6(3# )), the proportion of shear fracture reduces. In addition, the second phase particles were clearly found in dimples. Fig. 7 shows the SEM images of second phase particles in dimples of alloy 3# . The results, in conjunction with XRD in Fig. 3(d), confirm that the particles are Al3 CuGd containing Y. Two main reasons which can lead to the improvement of tensile strength by Gd, Y additions in the T6-tempered Al–Zn–Mg–Cu–Zr alloys are: on the one hand, a larger volume fraction and a smaller size of incoherent Al3 CuGd precipitated phases provide Orowan and order strengthening; on the other hand, inhibiting recrystallization and stabilizing deformation reversion microstructure during extrusion and solution treatments are notable and provide substructure strengthening. With these two comprehensive effects, the strength of the alloys are improved significantly. The improvement of the material ductility is largely due to the Al3 CuGd phase are uniform and nanoscale. These particles can effectively pin dislocations and subgrain boundary, inhibiting further recrystallization and grain growth during heat treatment and deformation process of material and having the fibrous unrecrystallized microstructure that make the alloy have good toughness. 4. Conclusions 1) After adding Gd or both with Gd and Y, the microstructures of as-cast A1–Zn–Mg–Cu–Zr alloys are significantly refined. Compared with non-RE addition, the alloys will form dispersed Al3 CuGd phase, which can strongly pin dislocations and subgrain boundary, inhibiting further recrystallization. The inhibition
recrystallization effect is more obvious when adding complex Gd and Y. 2) The micro-segregation of the Al–Zn–Mg–Cu–Zr alloy can be significantly reduced after adding Gd only or both with Gd and Y. Meanwhile, the grain relative solid-solubility of Zn, Mg and Cu elements could be improved, that is: from 54.82%, 49.63%, 26.38% to 69.78%, 78.75%, 42.19% after adding 0.25%Gd, and to 69.57%, 79.15%, 49.73% with 0.2% combined Gd and Y. 3) It is similar to the effect of artificial aging that adding single Gd or composite Gd and Y, which both can accelerate the peak aging time, improve the peak hardness, and extend the duration of peak aging time. 4) Compared with Al–Zn–Mg–Cu–Zr alloy, the alloys with Gd or with Gd and Y can significantly improve the strength and elongation of T6 state. When adding compound Gd and Y elements, the best mechanical properties are tensile strength is 639.3 Mpa, yield strength is 623.8 Mpa and elongation is 16.8% separately. Acknowledgment This research is financially supported by the National Natural Science Foundation of China (no. 50875031) and the Open Project from State Key Laboratory for Die Arrangement Technology of Huazhong University of Science and Technology (no. 136091320). References [1] M. Kevin, 3rd Annual AMM Aerospace Metals Conference, 2009. [2] V.V. Antipov, O.G. Senatorova, E.A. Tkachenko, R.O. Vakhromov, Met. Sci. Heat Treat. 9 (2011) 27–33. [3] W. Cassada, J. Liu, J. Staley, Adv. Mater. Proc. 160 (12) (2002) 27. [4] S. Bai, Z.Y. Liu, Y.T. Li, et al., Mater. Sci. Eng. A 527 (7–8) (2010) 1806–1814. [5] J. Zhong, Z. He, Aerospace Mater. 12 (2006) 1–13. [6] S.D. Liu, X.M. Zhang, M.A. Chen, Trans. Nonferrous Met. Soc. China 17 (4) (2007) 787–792. [7] S.H. Seyed Ebrahimi, M. Emamy, N. Pourkia, H.R. Lashgari, Mater. Des. 31 (2010) 4450–4456. [8] H. Li, Y. Yang, Z. Zheng, et al., Mater. Sci. Technol. 14 (1) (2006) 46–49. [9] O.N. Senkov, M.R. Shagiev, S.V. Senkova, Acta Mater. 56 (15) (2008) 3723–3738. [10] J. Yang, Z. Nie, T. Jin, et al., Chin. J. Nonferrous Met. l4 (4) (2004) 620–626. [11] X. Wang, Z. Nie, Spec. Cast. Nonferrous Alloys 29 (1) (2009) 76–79. [12] Z. Chen, X. Zhou, J. Shu, et al., Min. Metall. Eng. 2 (30) (2010) 102–106. [13] A.K. Chaubey, S. Mohapatra, B. Bhoi, et al., Defect Diffus. Forum 279 (2008) 97–103. [14] Q. Wang, D. Wang, J. China Univ. Min. Technol. 28 (4) (1999) 382–385. [15] S. Wang, S. Yang, C. Fang, et al., Chin. J. Nonferrous Met. 19 (12) (2009) 2083–2089. [16] G. Hu, X. Cai, Materials Science Foundation, vol. 9, Shanghai Jiao Tong University Press, Shanghai, 2007, p. 290. [17] A.C. Fionnuala, D.R. Joseph, B.P. Prangnell, Mater. Sci. Forum 757 (2002) 396–402. [18] L.F. Mondolfo, Aluminum Alloys: Structure and Properties, Butterworths, London, 1976. [19] V. Sudarshan, Trans. Indian Institute Met. 62 (3) (2009) 209–222.
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effects of Gd and Y additions on microstructure and properties of Al–Zn–Mg–Cu–Zr alloys XingGuo Zhang ∗ , FeiQiang Mei, HuanYue Zhang, ShaoHua Wang, CanFeng Fang, Hai Hao Foundry Center, Dalian University of Technology, Dalian 116024, China
a r t i c l e
i n f o
Article history: Received 29 March 2012 Received in revised form 11 May 2012 Accepted 13 May 2012 Available online 19 May 2012 Keywords: Gd Y Microstructure Mechanical properties
a b s t r a c t Three kinds of alloys Al–Zn–Mg–Cu–Zr, A1–Zn–Mg–Cu–Zr–Gd and A1–Zn–Mg–Cu–Zr–Gd–Y were prepared by cast metallurgy. The effects of Gd and Y additions on the microstructure and properties of Al–Zr–Mg–Cu–Zr alloys were investigated by optical microscopy (OM), tensile test, scanning electron microscopy (SEM), electronic probe microanalysis (EPMA) and transmission electronic microscopy (TEM). The results show that the combined additions of Gd and Y to the A1–Zn–Mg–Cu–Zr alloys effectively retard the mergence and growth of the sub-grain during the recovery process, significantly inhibit the recrystallization of Al matrix. b , 0.2 and ı are improved with increasing Gd and Y content, and the improving amplitudes reach 11.49%, 14.61%, 63.11% respectively and have better relative intracrystalline solubility. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Super-high strength aluminum alloy, which was set against the background of aerospace materials in 1960s, is a kind of high-strength aluminum alloy materials [1–3]. High strength, low density and good mechanical properties make Al–Zn–Mg–Cu alloys attractive for applications in the military and aerospace industries. But with the rapid development of composite material and titanium alloy, ultra-high strength aluminum alloy is facing unprecedented challenges. Improving the general properties of ultra-high strength aluminum alloy has been more and more focusing on microalloying. It was well accepted that trace additions of rare earth elements are of great importance to improve the microstructure and mechanical properties of aluminum alloys [4,5]. Zr, which can improve the strength, fracture toughness and stress corrosion resistance performance of ultra-high strength aluminum alloys, now has become the essential element for Al–Zn–Mg–Cu alloy [6,7]. Sc has been studied mostly as beneficial micro alloying element due to the presence of elastically hard, coherent and nano-sized Al3 Sc particles, which strongly refine as-cast grain size, inhibit the dynamic recrystallization and dislocation movement [8]. Furthermore, combined additions of Sc and Zr are shown to be more effective in refining as-cast grain size and improving mechanical properties because of the newly formed Al3 (Sc, Zr) particles in the Al–Zn–Mg–Cu alloys [9]. Y can effectively refine the branch crystal
∗ Corresponding author. Tel.: +86 0411 84706183; fax: +86 0411 84706183. E-mail address: [email protected] (X. Zhang). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.05.035
microstructure of as-cast alloy, inhibit the formation of bulky eutectic microstructure, and improve the elongation [10,11]. The micro alloying researches of rare earth among the aluminium alloy have developed from single microalloying to composite microalloying. But due to the complexity of the aluminium alloy systems as well as the limit and difficulty of characterizing the effect of rare earth. What is more, there are no complete theories available to guide the design of this type due to the lack of research on it. Setting up the relevant model of rare-earth micro alloying as soon as possible provides a new way to predict and improve the final properties [12]. Under this background, this paper will discuss the effect of the composite microalloying of the rare earth Gd and Y in Al–Zn–Mg–Cu–Zr super-high aluminium alloy, and provide reference to rare-earth micro alloying through cast ingot metallurgy process.
2. Experimental 2.1. Materials and processing The alloys investigated belong to the ultra-high strength Al–Zn–Mg–Cu–Zr alloys and include three kinds of alloys: Al–Zn–Mg–Cu–Zr, Al–Zn–Mg–Cu–Zr alloy with Gd additions (Al–Zn–Mg–Cu–Zr–0.25Gd), Al–Zn–Mg–Cu–Zr alloy with Gd and Y additions (Al–Zn–Mg–Cu–Zr–0.1Gd–0.1Y). The chemical composition of the alloys, analyzed by XRF-1800 type of X-ray fluorescence spectrum analyzer, is given in Table 1. High purity Al (99.99%), Mg (99.9%) and Gd (99.9%), and Al–50%Cu, Al–4%Zr, Mg–40%Y master alloys were taken as the starting raw material. The alloys were
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Table 1 Chemical composition of Al–Zn–Mg–Cu–Zr–Gd–Y alloys (mass fraction, %). Alloy
Zn
Mg
Cu
Zr
Gd
Y
1# 2# 3#
7.36 7.27 7.18
1.68 1.47 1.55
1.66 1.71 1.53
0.14 0.12 0.16
0 0.25 0.13
0 0 0.07
melted in an electrical-resistance furnace and cast into an iron mould.
2.2. Heat treatments The as-cast ingots were homogenized at 400 ◦ C for 4 h plus 460 ◦ C for 24 h, followed by air cooling to the room temperature. Then the ingots were converted into plates, with an rolling ratio of 10:1. Subsequently, the samples were solid solution treated at 470 ◦ C for 2 h and quenched in ambient temperature water immediately (no stretching was applied). Following quenching, the plates were aged at 120 ◦ C for 24 h for T6 tempered.
Fe ≤ 0.05 ≤ 0.05 ≤ 0.05
Si ≤ 0.05 ≤ 0.05 ≤ 0.05
Al Bal. Bal. Bal.
2.4. Microstructural characterizations The microstructure was observed by MEFS multifunction optical microscopy. The fractures were observed by scanning electro microscope (SEM). The thin foils for transmission electron microscopy (TEM) were prepared by twin jet-polishing in electrolyte solution of HNO3 and methyl alcohol (1:3). TEM examinations were performed through a JEM-2010 microscope operated at 200 kV. X-ray diffraction analysis (XRD) was used to characterize the phases of samples on a D/max-2500 diffractometer with Cu K␣ radiation. 3. Results and discussion
2.3. Mechanical properties and relative intracrystalline solubility
3.1. Microstructure of the alloys
The tensile properties were tested at room temperature by an electro universal test machine. The tests were performed at a nominal tension rate of 2 mm/min according to the IGC test standard of GB/T 228-2002. The hardness that can be used to monitor the aging processing was measured by light loaded durometer. The magnitude of load was 1000 kg with a 30 s load time. The relative intracrystalline solubility of main alloy elements in the grain was determined by electron probe microanalysis (EPMA) with EPMA-1600.
The microstructures of the studied alloys are shown in Fig. 1. Coarsening as-cast grains (the average size is 62 m), dendritic structures and serious segregation can be obviously observed in alloy 1# . Adding rare-earth could greatly refine the as-cast grain and abated the dendritic structures. The most effective refining appears in alloy 2# , with the average size 32 m, and most dendritic structures vanish (as shown in Fig. 1(2# )). Compared to alloy 2# , simultaneously adding of Gd and Y, the microstructure appears a certain coarsening, with the average size 40 m, but is still basically equiaxial crystal.
Fig. 1. Optical micrographs of the as-cast microstructure.
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Fig. 2. Optical microstructures of two studied alloys with solution treatment.
With the adding of Gd and Y, crystalline grains of ␣(Al) of the alloys in the as-cast structure are refined, that is related to the characteristic of Gd and Y in the process of solidification of the alloys. The chemical properties of Gd and Y elements are rather active and the electronegativities of Gd and Y are 1.2 and 1.22, respectively, smaller than Al(1.5). They can fill the vacancies on the surfaces when they dissolve in the alloys which can reduce the surface tension between two phases. Consequently, nucleation speed can increase. Meanwhile, active surface membranes form between grains and the liquid alloy which block the growing of the grains and make the grains refined. Grain refinement can be obtained that compared with Al, the atomic radius of Gd and Y are larger. The results show the substitutional solid solution is formed when Gd and Y additions are dissolved into Al [13,14]. When Y is added into Al–Zn–Mg–Cu–Zr–Gd alloy, the effect of grain refinement is weaker than that adding Gd element only. It is mainly due to the interaction among Zr, Gd and Y elements that weaken the effect of grain refinement [11]. Typical optical microstructures of T6-tempered Al–Zn–Mg–Cu–Zr–Gd–Y alloys, which had been rolled, solution strengthened, T6 ageing strengthened and corroded by Kelvin solution, are shown in Fig. 2. We could conclude from the pictures that Al–Zn–Mg–Cu–Zr alloy is recrystallized. Some of the subgrains grow up into large recrystallized grains and some of the subgrains boundary disappear
(Fig. 2(1# )). Al–Zn–Mg–Cu–Zr–0.25Gd alloy is partially recrystallized (Fig. 2(2# )). Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy has a fibrous unrecrystallized microstructure with fine subgrain microstructures (Fig. 2(3# )). It is clear that the recrystallization resistance of T6-tempered Al–Zn–Mg–Cu–Zr alloy can be further enhanced by adding Gd and Y in combination. TEM micrographs and XRD of T6-tempered The Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy are presented in Fig. 3. The 10–20 nm spherical dispersoids of second phase have precipitated from matrix after 120 ◦ C, 24 h aging treatment (as shown in Fig. 3(a) and (b)). The chemical composition of the dispersoids was determined by EDXS. The spherical dispersoids contain Gd, Y, Zn, Cu and Al. Based on the XRD (as shown in Fig. 3(d)) and EDXS, it is suggested that Y was dissolved in Al3 CuGd. The Al3 CuGd particles which can strongly pin dislocations and grain boundary (as shown in Fig. 3(c)) appear in the alloys. Thus these particles block the movement of dislocations as well as subboundary and have strong stabilization effect of the substructure of the deformation organization. 3.2. Relative intracrystalline solubility The EPMA images and distribution of major elements of alloys in 1# , 2# and 3# are shown in Fig. 4. The EPMA image of alloy 1# shows that major elements (Zn, Mg and Cu) in Al–Zn–Mg–Cu–Zr
Fig. 3. The TEM micrographs and XRD of T6-tempered Al–Zn–Mg–Cu–Zr–0.2% (Gd + Y) alloy.
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Fig. 4. Mapping analyses of main alloy elements in Al–Zn–Mg–Cu–Zr–Gd–Y alloy billets.
enrich on the boundary of the grain (as shown in Fig. 4(1# )). These results indicate that the concentration of elements on the boundary is higher than that in the grain, especially Cu. This is because during the solidification process of alloys, Cu is vulnerable to segregate in the grain boundary in the form of non-equilibrium eutectic compound or the insoluble second phase particles. In the specimen with no additional rare-earth elements, Zn and Mg are rarely dissolved in the grain, but most of them are eutectic segregation on grain boundary. In alloy 2# and 3# , although there is a large concentration gradient of Cu, the segregation of alloying element Zn and Mg on the boundary reduces which increases the grain relative solid-solubility and reduces the micro-segregation of the alloys. The segregation degree of solute elements of the ingot casting is suppressed greatly. In order to point out the grain relative solid-solubility of major elements and offer reference for the following heat treatment technology, this paper adopt the relative intracrystalline solubility of elements (the relative intracrystalline solubility in the grain of element A = content of element A in grain/content of alloying element A) to represent the effect of different contents of rare earth content on the solid-solubility [15]. The results are shown in Table 2. The micro-segregation of the Al–Zn–Mg–Cu–Zr alloy could be significantly reduced after adding Gd or both with Gd and Y. Meanwhile, the relative intracrystalline solubility of Zn, Mg and Cu elements could be improved, that is: from 54.82%, 49.63%, 26.38% to 69.78%,
78.75%, 42.19% after adding 0.25% Gd, and to 69.57%, 79.15%, 49.73% after adding 0.2% combined Gd and Y respectively. According to modern statistical physics, equilibrium partition coefficient can be expressed in diffusion activation energy [16]: k0 = exp
(Q − Q ) L S KT
where QS is diffusion activation energy which the solute atoms needed to across the interface barrier layer into the melt; QL is diffusion activation energy which the solute atoms needed to across the interface barrier layer into the solid solution; K is Boltzmann constant; T is phase transition temperature. Distribution coefficient of solute essentially reflects the difference that solute atoms through the solid-liquid boundary layer to the sides on micro-interface boundary. The equilibrium partition coefficient of Gd and Y are less than 1 and the solid solubility in the Table 2 Relative intracrystalline solubility in the grain of main alloy elements. Alloy
1# 2# 3#
Relative intracrytalline solubility Zn
Mg
Cu
54.82 69.78 69.57
49.63 78.75 79.15
26.38 42.19 49.73
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Fig. 5. Ageing–hardening curves of artificial ageing of alloys.
aluminum is very small, so most of them will enrich in the liquid boundary, which reduces the chance of Zn, Mg and Cu moving into the solid solution and also increases the solid solubility of Zn, Mg and Cu in grain, reduces the microsegregation of alloy. 3.3. Age-hardening curves of the alloys The artificial age-hardening curves of alloys are shown in Fig. 5. It indicates that the aging strengthening properties of these studied alloys are obviously improved with prolonging aging time. The improvements of peak-aging hardness are all above 50% and the artificial age-hardening curves of the alloy 2# and alloy 3# stay consistent approximately. In the process of artificial age-hardening, the peak-aging of alloy 1# locates at 11–24 h while alloy 2# and alloy 3# appear after 7 h and last up to 32 h. Compared to the alloy 1# , the
peak-aging hardness of the alloy 2# and the alloy 3# improve about 20 HBS. The results show that the addition of rare earth extends the peak-aging time, makes the peak-aging arriving earlier and improves the peak-aging hardness of artificial age-hardening at 120 ◦ C. The mainly strengthened phase in Al–Zn–Mg(–Cu) alloy is (MgZn2 ) and , and the general precipitation sequence is as follows: a sosoloid → G.P. zones → (MgZn2 ) → phase. Two kinds of effects of Gd and Y can be observed in Al–Zn–Mg–Cu–Zr alloys. One is minor rare earth element addition can improve the concentration of Zn, Mg and Cu into a sosoloid, as suggested before, therefore the driving force of phase transition is enhanced from G.P. zones to (MgZn2 ) phase [17]. The other is Al3 CuGd particles can come into being nucleation core during aging when these particles do not melt back after solution treatment and distribute dispersively in the alloy matrix. Thus minor rare earth element addition can make the peak aging arriving earlier for Al–Zn–Mg–Cu–Zr alloys. Gd and Y are elements that display slightly lower binding energy with vacancies than are exhibited by Mg and Zn [18,19], which will automatically preferentially attach to vacancies and sequester them. Such a preferential binding of vacancies with Gd and Y or both results in a decrease in the number of vacancies available for the movement of Mg and Zn atoms and hence is exhibited in delayed growth of the ’ (MgZn2 ) phase. Thus the observed delay in reaching the peak after adding Gd only or both with Gd and Y is directly related to the binding energy effect. 3.4. Tensile properties of the alloys under T6 treatment The tensile properties of the alloys under T6 treatment are shown in Table 3. It indicates that compared with non-RE addition alloy, adding Gd only or both with Gd and Y can significantly improve the strength and elongation of T6 state. When the alloy containing 0.25% Gd, the ultimate tensile strength and yield strength values are 624.54 MPa, 595.00 MPa at T6 and the
Fig. 6. SEM micrographs of tensile fracture surface of alloys.
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Fig. 7. Morphology and composition of second-phase in the dimple.
Table 3 Mechanical properties of alloys. Alloy #
1 2# 3#
b (MPa)
0.2 (MPa)
ı (%)
573.44 624.54 639.30
544.30 595.00 623.80
10.3 13.3 16.8
elongation is 13.3%. The improving amplitudes reach 8.91%, 9.31%, 29.13%, respectively. When adding both compound Gd and Y elements, the properties are optimal: tensile strength is up to 639.3 Mpa, yield strength is 623.8 Mpa, elongation is 16.8%, and the improving amplitudes reach 11.49%, 14.61%, 63.11%, respectively. Fig. 6 shows the SEM fracture morphology of the tensile test specimen under T6 treatment. The main fracture behavior of Al–Zn–Mg–Cu–Zr alloy with no additional rare-earth elements is shear fracture and dimple intercrystalline fracture (as shown in Fig. 6(1# )), whereas the primary fracture mechanism of the alloy with additional rare-earth elements is dimpled transgranular type of failure (as shown in Fig. 6(2# ), Fig. 6(3# )), the proportion of shear fracture reduces. In addition, the second phase particles were clearly found in dimples. Fig. 7 shows the SEM images of second phase particles in dimples of alloy 3# . The results, in conjunction with XRD in Fig. 3(d), confirm that the particles are Al3 CuGd containing Y. Two main reasons which can lead to the improvement of tensile strength by Gd, Y additions in the T6-tempered Al–Zn–Mg–Cu–Zr alloys are: on the one hand, a larger volume fraction and a smaller size of incoherent Al3 CuGd precipitated phases provide Orowan and order strengthening; on the other hand, inhibiting recrystallization and stabilizing deformation reversion microstructure during extrusion and solution treatments are notable and provide substructure strengthening. With these two comprehensive effects, the strength of the alloys are improved significantly. The improvement of the material ductility is largely due to the Al3 CuGd phase are uniform and nanoscale. These particles can effectively pin dislocations and subgrain boundary, inhibiting further recrystallization and grain growth during heat treatment and deformation process of material and having the fibrous unrecrystallized microstructure that make the alloy have good toughness. 4. Conclusions 1) After adding Gd or both with Gd and Y, the microstructures of as-cast A1–Zn–Mg–Cu–Zr alloys are significantly refined. Compared with non-RE addition, the alloys will form dispersed Al3 CuGd phase, which can strongly pin dislocations and subgrain boundary, inhibiting further recrystallization. The inhibition
recrystallization effect is more obvious when adding complex Gd and Y. 2) The micro-segregation of the Al–Zn–Mg–Cu–Zr alloy can be significantly reduced after adding Gd only or both with Gd and Y. Meanwhile, the grain relative solid-solubility of Zn, Mg and Cu elements could be improved, that is: from 54.82%, 49.63%, 26.38% to 69.78%, 78.75%, 42.19% after adding 0.25%Gd, and to 69.57%, 79.15%, 49.73% with 0.2% combined Gd and Y. 3) It is similar to the effect of artificial aging that adding single Gd or composite Gd and Y, which both can accelerate the peak aging time, improve the peak hardness, and extend the duration of peak aging time. 4) Compared with Al–Zn–Mg–Cu–Zr alloy, the alloys with Gd or with Gd and Y can significantly improve the strength and elongation of T6 state. When adding compound Gd and Y elements, the best mechanical properties are tensile strength is 639.3 Mpa, yield strength is 623.8 Mpa and elongation is 16.8% separately. Acknowledgment This research is financially supported by the National Natural Science Foundation of China (no. 50875031) and the Open Project from State Key Laboratory for Die Arrangement Technology of Huazhong University of Science and Technology (no. 136091320). References [1] M. Kevin, 3rd Annual AMM Aerospace Metals Conference, 2009. [2] V.V. Antipov, O.G. Senatorova, E.A. Tkachenko, R.O. Vakhromov, Met. Sci. Heat Treat. 9 (2011) 27–33. [3] W. Cassada, J. Liu, J. Staley, Adv. Mater. Proc. 160 (12) (2002) 27. [4] S. Bai, Z.Y. Liu, Y.T. Li, et al., Mater. Sci. Eng. A 527 (7–8) (2010) 1806–1814. [5] J. Zhong, Z. He, Aerospace Mater. 12 (2006) 1–13. [6] S.D. Liu, X.M. Zhang, M.A. Chen, Trans. Nonferrous Met. Soc. China 17 (4) (2007) 787–792. [7] S.H. Seyed Ebrahimi, M. Emamy, N. Pourkia, H.R. Lashgari, Mater. Des. 31 (2010) 4450–4456. [8] H. Li, Y. Yang, Z. Zheng, et al., Mater. Sci. Technol. 14 (1) (2006) 46–49. [9] O.N. Senkov, M.R. Shagiev, S.V. Senkova, Acta Mater. 56 (15) (2008) 3723–3738. [10] J. Yang, Z. Nie, T. Jin, et al., Chin. J. Nonferrous Met. l4 (4) (2004) 620–626. [11] X. Wang, Z. Nie, Spec. Cast. Nonferrous Alloys 29 (1) (2009) 76–79. [12] Z. Chen, X. Zhou, J. Shu, et al., Min. Metall. Eng. 2 (30) (2010) 102–106. [13] A.K. Chaubey, S. Mohapatra, B. Bhoi, et al., Defect Diffus. Forum 279 (2008) 97–103. [14] Q. Wang, D. Wang, J. China Univ. Min. Technol. 28 (4) (1999) 382–385. [15] S. Wang, S. Yang, C. Fang, et al., Chin. J. Nonferrous Met. 19 (12) (2009) 2083–2089. [16] G. Hu, X. Cai, Materials Science Foundation, vol. 9, Shanghai Jiao Tong University Press, Shanghai, 2007, p. 290. [17] A.C. Fionnuala, D.R. Joseph, B.P. Prangnell, Mater. Sci. Forum 757 (2002) 396–402. [18] L.F. Mondolfo, Aluminum Alloys: Structure and Properties, Butterworths, London, 1976. [19] V. Sudarshan, Trans. Indian Institute Met. 62 (3) (2009) 209–222.