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Effect of Sc and Zr additions on microstructures and corrosion behavior of Al-Cu-Mg-Sc-Zr alloys
Accepted Manuscript Title: Effect of Sc and Zr additions on microstructures and corrosion behavior of Al-Cu-Mg-Sc-Zr alloys Author: Fangfang Sun Guiru Liu Nash Qunying Li Enzuo Liu Chunnain He Chunsheng Shi Naiqin Zhao PII: DOI: Reference:
S1005-0302(17)30004-X http://dx.doi.org/doi:10.1016/j.jmst.2016.12.003 JMST 883
To appear in: Received date: Revised date: Accepted date:
6-9-2016 31-10-2016 19-12-2016
Please cite this article as: {http://dx.doi.org/
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Effect of Sc and Zr Additions on Microstructures and Corrosion Behavior of Al-Cu-Mg-Sc-Zr Alloys Fangfang Sun1, Guiru Liu Nash2, Qunying Li1, Enzuo Liu1,3, Chunnain He1,3, Chunsheng Shi1,3, Naiqin Zhao1,3* 1
School of Materials Science and Engineering, Tianjin University, Tianjin 300072, China
2
Department of Mechanical, Materials and Aerospace Engineering, Illinois Institute of Technology,
USA 3
Tianjin Key Laboratory of Composite and Functional Materials, Tianjin 300072, China
*Corresponding author. Tel. & Fax: +86 1 27891371. E-mail address: [email protected] (N. Zhao).
[Received 6 September 2016; Received in revised form 31 October 2016; Accepted 19 December 2016]
Graphical Abstract
1
Highlights: 1.
Adding Sc and Zr changed the cross section morphology of inter-granular corrosion.
2.
There exists a corrosion mechanism conversion of S precipitates in Al-Cu-Mg alloy.
3.
Adding Sc and Zr restricted the corrosion mechanism conversion of S particles.
4.
The Al-Cu-Mg-Sc-Zr alloy possesses excellent inter-granular corrosion properties.
5.
The content of Cu and Mg elements decreased along the grain boundary after minor Sc and Zr addition.
2
The effects of adding the alloy element Sc to Al alloys on strengthening, recrystallization and modification of the grain microstructure have been investigated. The combination of Sc and Zr alloying not only produces a remarkable synergistic effect of inhibition of recrystallization and refinement of grain size but also substantially reduce the amount of high-cost additional Sc. In this work, the microstructures and corrosion behavior of a new type of Al-Cu-Mg-Sc-Zr alloy with Sc/Zr ratio of 1/2 were investigated. The experimental results showed that the Sc and Zr additions to Al-Cu-Mg alloy could strongly inhibit recrystallization, refine grain size, impede the segregation of Cu element along the grain boundary and increase the spacing of grain boundary precipitates. In addition, adding Sc and Zr to Al-Cu-Mg alloy effectively restricts the corrosion mechanism conversion associated with Al2CuMg particles, which resulted in the change of the cross-section morphology of inter-granular corrosion from an undercutting to an elliptical shape. The susceptibility to inter-granular corrosion was significantly decreased with increasing Sc and Zr additions to the Al-Cu-Mg alloy. The relationships between microstructures evolution and inter-granular corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys were also discussed.
Key words: Aluminum alloy; Microstructure; Scandium; Inter-granular corrosion
1. Introduction Al-Cu-Mg alloys are important structural materials that have been widely used in the aerospace field due to the high specific strength, excellent heat resistance and easy processability[1-3]. However, the tendency for inter-granular corrosion (IGC) of Al-Cu-Mg alloys is relatively high due to the potential difference between the heterogeneous distributions of S (Al2CuMg) or θ (Al2Cu) precipitates along the grain boundaries and in the matrix. Therefore, development of high specific strength with excellent corrosion resistance of Al-Cu-Mg alloys is a main research goal for potential applications in aviation, aerospace, and other high-technology sectors[4]. Micro-alloying with rare earth or transition metals (such as Sc, Zr, Cr, Nb, Sr and Ti) was discovered as an important method to modify microstructures and improve performance of the aluminum alloy[5-11]. Compared to the Al-4Mg alloy, the addition of Sc and Yb reduced the grain size, 3
changed the morphology of intermetallic particles and decreased the corrosion potential values of Al-4Mg alloy in 3.5 wt% NaCl solution[12]. Gupta et al.[13,14] studied the sensitization behavior of Al-5Mg alloy using nitric acid mass loss test found that micro-alloying additions Nd ( >0.11%), Sr, Si, and Ti significantly reduced IGC susceptibility of Al-5Mg alloy. Moreover, Fang et al.[15] and Peng et al.[16] showed that complex additions of Er, Zr and Cr to Al-Zn-Mg-Cu alloys remarkably enhanced resistance to IGC, which was contributed by the formation of 20–50 nm particles strongly inhibiting the recrystallization behavior of the Al matrix. However, Sc was reported to be the most effective micro-alloying element to modify microstructures and improve properties of Al alloy[17]. The quench sensitivity of the Al-7Si-0.6 Mg alloy was effectively improved by more than 60% with 0.04% Sc additions[18]. The effects of Sc additions on microstructure and properties of rheo-diecasting wrought aluminum alloy were investigated by Zhao et al.[19]. This study concluded that Sc additions not only influenced the morphology of the primary α-Al particles, but also promoted the formation of fine α-Al globules during the solidification process of the remaining liquid, resulting in the considerable improvement in mechanical properties. Moreover, Sc-containing Al alloys are limited by high cost. Previous studies had demonstrated that the combined alloying elements of Sc and Zr not only produced remarkable synergistic effects but also substantially reduced the Sc content. Qian et al.[20] systematically investigated the mechanical properties of L12-Al3(Sc,Zr) using first principles calculations. The results showed that the ideal strength and ductility of L12- Al3(Sc,Zr) increased with increasing Zr addition. Li et al.[21] investigated different Sc/Zr ratio addition to the Al-Zn-Mg-Cu alloy and concluded that the strengthening effect was better with a Sc/Zr ratio of 2/4 than that without Sc and Zr additions or only adding Sc. Most of the literature has focused on microstructure and mechanical properties, while studies on corrosion behavior of Sc- and Zr-containing Al alloys, especially of Sc- and Zr-containing Al-Cu-Mg alloys, are scarce[5,12,21-25]. The improvements of mechanical properties of Al-7Si-0.65Mg alloy were attributed to microstructural refinement, particularly the modification of eutectic Si and precipitation of secondary nano-scale Al3(Sc, Zr) dispersoids after minor Sc and Zr additions[22]. The effect of Sc and Zr on corrosion behavior of Al-Mg alloys[9, 12] and Al-Zn-Mg alloys[4, 6, 15, 26] have been reported. Li et al.[26] pointed out that the enhancement of corrosion performance of Al-Zn-Mg-Mn alloy was due to the effect of inhibiting recrystallization behavior caused by Al3 (Sc, Zr) precipitates and interrupted distributions of grain boundary precipitates. Little research has been made on effects of 4
Sc and Zr alloying element additions on microstructure and corrosion properties of Al-Cu-Mg alloys. Furthermore, Zhu et al.[27] deemed that AA2024 (Al-Cu-Mg alloys) displayed a completely different morphology of corrosion cross-section from 7B04 aluminum alloys. The relationship between the corrosion property, microstructure and the cross-section morphology of the inter-granular corrosion with Sc and Zr alloying element additions to Al-Cu-Mg alloys is still unclear. Therefore, it is necessary to study the relationship between corrosion resistance (including cross-section morphology of inter-granular corrosion) and micro-alloying element addition to understand the corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys. The purpose of this work is to conduct a comprehensive study the effect of a specific Sc/Zr ratio on microstructure and corrosion behavior of Al-Cu-Mg alloys. The microstructure evolution, cross-section morphology of inter-granular corrosion and corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys will be discussed in this paper. 2. Materials and Experimental Methods 2.1 Materials preparation Three types of Al-Cu-Mg (in wt%) alloys with Sc and Zr additions at a constant Sc/Zr ratio of 1/2 were prepared. In order to fully nucleate Al3(Sc, Zr) particles, we designed the alloy with a slight excess Zr content, since the Al3Sc phase nucleates on the surface of the Al3Zr phase particle. Flat bar ingots with 20 mm in thickness, were provided by Hunan Rare-Earth Metal Research Institute, China. The chemical compositions of the alloys were analyzed by inductively coupled plasma–optical emission spectrometry as shown in Table 1. The ingots were homogenized at 480 oC for 20 h, hot rolled from 20 to 6.9 mm and then cold rolled to sheets with a 2.7 mm final thickness. The total deformation reduction of each alloy was 86.5%. The sheets of all three alloys were solution treated at 500 oC for 1 h, subsequently water quenched and artificially aged at 190 oC for various time. Ten hardness measurements were conducted on each sample and the average results with the standard deviation from the mean were reported. Aging hardening curves of the Al-Cu-Mg(-Sc-Zr) alloys at 190 oC are presented in Fig. 1. It was notable that all of the standard deviations of the samples were below HV3.5. The hardening behavior of the alloys aged at 190 oC displayed a dual-peak phenomenon. After minor Sc and Zr addition, the second peak-aged time was reduced and the peak hardness was increased with increasing amounts of Sc and Zr. The peak-ageing time of Al-Cu-Mg, Al-Cu-Mg-0.3(Sc+Zr) and Al-Cu-Mg-0.6(Sc+Zr) alloys were 22, 10, and 10 h, respectively. 5
Hereafter, all of the alloys are in the peak-aged condition unless otherwise specified. 2.2 Microstructural characterization Metallographic observation was conducted for the solid–solution state samples using a Leica optical microscope. The samples were prepared by conventional mechanical polishing and then etched with Keller's reagent (95 vol.% H2O + 2.5 vol.% HNO3 + 1.5 vol.% HCl + 1 vol.% HF). Misorientation angles between grains/sub-grains of T6 treated specimens were measured by electron back scattered diffraction (EBSD). The microstructure of peak-aged sheets was analyzed by transmission electron microscopy (TEM). Energy dispersive spectroscopy (EDS) was utilized to investigate the composition of grain boundary precipitates. Thin foils for TEM observations were prepared by twin-jet electro-polishing at 42.5 mA in a solution of 30% nitric acid and 70% methanol solution cooled to -30 oC. A JEM-2100F electron microscope and a TECNAI G2 F20 electron microscope were used to conduct the analysis. 2.3 Corrosion tests After grinding and polishing, all the peak-aged specimens were immersed in the corrosive solution (57 g NaCl + 1 l H2O + 10 ml H2O2) at (35±2) oC for 6 h. The ratio of the exposed area of specimens and the volume of corrosion solution was less than 20 mm2/ml, according to inter-granular corrosion testing requirement of GB/T 7998-2005[28]. In order to evaluate the mechanical properties, tensile tests were conducted using the samples with and without immersing in the corrosive environment. Tensile tests were conducted using a CSS-44100 electronic universal testing machine with 1 mm/min loading speed. The axis of the tensile specimens was parallel to the rolling direction of the thin sheets. All experiments were conducted three times and the average testing results are reported with the standard deviation. The cross-section (perpendicular to the rolling direction) of samples was measured using a Leica optical microscope after the inter-granular corrosion tests. After removing the corrosion products from the test specimens, energy dispersive spectroscopy (SEM-EDS) was utilized to investigate the composition of the region at the bottom of inter-granular corrosion with prolonged time for immersion in the corrosive solution. 3. Results 3.1
Microstructures Fig. 2 shows the microstructures of three different alloys after solution treatment at 500 oC for 1
h. It was found that the grain size was reduced with increasing amounts of Sc and Zr in the 6
Al-Cu-Mg alloy. However, it showed that Al-Cu-Mg-0.1Sc-0.2Zr alloy had slightly smaller banded structure than Al-Cu-Mg-0.2Sc-0.4Zr alloy. Fig. 3 shows the SEM-EBSD images and misorientation angle distribution of three alloys. It was found that the Al-Cu-Mg alloy had uniaxial grains and the alloys with additional Sc and Zr had elongated grains as shown in Fig. 3(a, b and c). The boundary misorientation angle distributions of Al-Cu-Mg and Al-Cu-Mg-Sc-Zr alloys after peak ageing (190 o
C) are shown in Fig. 3(d, e and f). After the solution treatment at 500 oC for 1 h, full
recrystallization occurred in the Al-Cu-Mg alloy resulting in high angle grain boundaries (HAGBs) with the misorientation angles mainly ranging from 20° to 55° (Fig. 3(a and d)). However, Al-Cu-Mg alloy with Sc and Zr alloying element additions still had a fibrous unrecrystallized microstructure. The grain boundaries were characterized as low angle grain boundaries (LAGBs) with the misorientation angles mostly less than 10°. In addition, with the increase in the Sc and Zr alloying element additions, the proportion of LAGBs increased, indicating that the effect of inhibiting recrystallization was improved (Fig. 3(e and f)). The above observation was further verified by TEM images as shown in Fig. 4. It was found that the Al-Cu-Mg-Sc-Zr peak-aged alloys contained more micro-scale sub-grains than the recrystallized Al-Cu-Mg alloy as shown in Fig. 4(a, b and c). The grain boundary precipitates (GBPs) in Al-Cu-Mg alloy were distributed continuously, as shown in Fig. 4(a). The GBPs were distributed discontinuously by minor Sc and Zr additions as shown in the upper right corner of Fig. 4(b and c). In addition, the interparticle spacing of GBPs was increased with the increase in the Sc and Zr additions, as shown in Fig. 4(a, b and c). It was also found that the size of sub-grains decreased with the increase in the Sc and Zr alloying element additions as shown in Fig. 4(b and c). Very fine precipitate particles were found inside the sub-grains as shown in Fig. 4(d and e). The [001] selected area diffraction (SAD) pattern revealed that the fine precipitation particles were Al3(Sc,Zr) as shown in Fig. 4(f). The sub-grains might have resulted from the formation of nanometer sized, coherent, secondary Al3(Sc,Zr) particles with the additions of Sc and Zr. The [011] Al SAD pattern (Fig. 4(g)) revealed that light spots located in 1/3<200> and 2/3<200> directions corresponding to acicular exudation phase was S’ phase[29]. Thus, S’ phase is the main strengthening precipitate in the peak-aged condition of Al-Cu-Mg alloy. The chemical compositions of GBPs determined by TEM-EDS are listed in Table 2. It was found that the percentage of the Cu and Mg elements in GBPs were substantially decreased with the 7
additions of Sc and Zr, which implied that the segregation of Cu and Mg elements along the grain boundary was inhibited to some extent by minor Sc and Zr additions. This finding was consistent with previous results reported by Wang et al.[8], which supported the finding that the grain refinement was an effective method to reduce or eliminate grain boundary segregation. Furthermore, the value of the Cu/Mg ratio in the precipitates significantly decreased with increasing amount of Sc and Zr additions. 3.2 Inter-granular corrosion The cross section morphologies of inter-granular corrosion were analyzed by optical microscopy, as shown in Fig. 5. It was found in Fig. 5 that the maximum depth of inter-granular corrosion was 124 m for the Al-Cu-Mg alloy, 107 m for the Al-Cu-Mg-0.1Sc-0.2Zr alloy and 61m for the Al-Cu-Mg-0.2Sc-0.4Zr alloy. The dramatic decrease in the maximum depth of inter-granular corrosion indicated that the inter-granular corrosion resistance was greatly increased with minor Sc and Zr additions. In order to evaluate the susceptibility to inter-granular corrosion of the studied alloys, the ultimate tensile strength loss ratio (δUTS) and elongation loss ratio (δEI) were calculated as:
UTS
b - f b
(1)
EI
b - f b
(2)
where σb and δb are the ultimate tensile strength and total elongation of samples without immersion in the corrosive solution, respectively; σf and δf are the ultimate tensile strength and total elongation of the samples after immersion in the corrosive solution for 6 h at 35 oC. The variations of the δUTS and δEI with different amounts of Sc and Zr additions for the peak-aged Al-Cu-Mg alloys are listed in Table 3. It was found that the δUTS and δEI decreased with increasing amount of Sc and Zr additions, which indicated that the minor additions of Sc and Zr tended to decrease the inter-granular corrosion susceptibility of the Al-Cu-Mg alloys. The observation was consistent with the measured maximum depth of inter-granular corrosion as a function of the amount of Sc and Zr alloying element additions. Furthermore, the decrease in ultimate tensile strength of Al-Cu-Mg-0.2Sc-0.4Zr alloy might have resulted from the formation of coarse and primary Al3Zr particles when the content of Zr was over 0.3 wt%[30]. 8
3.3 Evolution of the cross section morphology of inter-granular corrosion The cross section morphology of the tested samples contains abundant characteristic information including corrosion types, corrosion mechanism, microstructure and residual fatigue strength[27]. A comprehensive investigation of the effect of minor Sc and Zr additions on the evolution of cross section morphology of inter-granular corrosion and corrosion mechanism has been carried out. Fig. 6 shows the cross section morphology of the Al-Cu-Mg and Al-Cu-Mg-0.2Sc-0.4Zr alloys after soaking in the inter-granular corrosion solution for different durations. Compared to Al-Cu-Mg alloy, it was found that corrosion area and corrosion depth for the same soaking duration were both significantly reduced by adding minor Sc and Zr elements. The cross section morphology of inter-granular corrosion of the Al-Cu-Mg alloy appeared to extend in a lateral direction below the surface. Therefore, the surface width of the cross section morphology in the Al-Cu-Mg alloy was small with a bigger inner width beneath the surface as shown in Fig. 6(a, c, e and g), which would have resulted in the miscalculation of the corrosion degree. However, the surface width of the cross-section morphology of inter-granular corrosion in the Al-Cu-Mg-0.2Sc-0.4Zr alloy was the largest and gradually reduced toward the bottom as shown in Fig. 6(b, d, f and h). A schematic of the cross section morphology of the two alloys are illustrated in Fig. 6(i and j). The overall topography of the cross section morphology is shown in Fig. 7. It was found that the edge of the cross section of inter-granular corrosion was smooth and bottom of cross section was uneven. Table 4 shows the chemical compositions at the bottoms of cross section of the Al-Cu-Mg and Al-Cu-Mg-0.2Sc-0.4Zr alloys, respectively. It was found that Cu was segregated at the bottom of the cross section in the Al-Cu-Mg alloy after a prolonged immersion in the corrosive solution. Compared to the Al-Cu-Mg alloy, the Cu and the Cu/Mg ratio were both significantly reduced at the bottom of cross section in the Al-Cu-Mg-0.2Sc-0.4Zr alloy for the same soaking time. 4. Discussion 4.1 Corrosion mechanism The combination of Sc and Zr significantly enhanced the resistance to recrystallization of the Al-Cu-Mg alloy, which can be attributed to the formation of high density, coherent and dispersed secondary Al3(Sc, Zr) particles. The Al3(Sc,Zr) particles were identified by both the presence of Ashby-Brown contrast (Fig. 4(d and e)) and the superstructure reflections like [001]Al (Fig. 4(f))[15]. 9
The coherent and fine dispersed secondary Al3(Sc,Zr) particles strongly pin dislocations and the sub-grain/grain boundaries during further heat treatment processes, resulting in unrecrystallized sub-grain structures retained in the deformed Al and Al alloys[5,21,31,32]. Vlach et al.[33] reported that combination of Sc and Zr additions to Al suppressed recrystallization during annealing at 550 oC. In our work, the effect of recrystallization inhibition was increased with increasing amount of Sc and Zr additions. It is well documented that the major microstructural features affecting the inter-granular corrosion resistance are the morphology and chemical composition of the GBPs[26, 34-37]. With the addition of Sc and Zr, the fiber-like sub-grains consisting of LAGBs were obtained. Compared to LAGBs, HAGBs have high energy, which is apt to form continuous and coarse precipitates along the grain boundaries. However, the LAGBs have low energy, which is prone to the formation of finer and discontinuous GBPs along the low angle unrecrystallized sub-grain boundaries. According to the electrochemical behavior of the precipitates in Al-Cu-Mg alloy, the potentials of S (S’, S’’) phase is more negative than the matrix and precipitate free zone (PFZ)[38]. Thus, S (S’, S’’) phase would act as the anode and firstly dissolve when the alloy is exposed to a corrosive environment according to the anodic dissolution mechanism. Therefore, the continuous S’ phase distribution in fully recrystallized Al-Cu-Mg alloy offers an unblocked path for corrosion. However, the discontinuous distribution of S’ precipitates along LAGBs could cut off the path of anode dissolution in the Al-Cu-Mg-Sc-Zr alloys, which results in the decrease in the corrosion speed and the improvement of corrosion resistance. In summary, adding Sc and Zr alloying elements to Al-Cu-Mg alloys can enhance the resistance to recrystallization and stabilize the deformed fibrous microstructure with numerous LAGBs due to the formation of Al3(Sc,Zr) particles. The deformed fibrous microstructure with the discontinuous precipitate distribution at the sub-grain boundaries resulted in the improvement in corrosion resistance. The more the Sc and Zr additions, the more remarkable the effect on improving corrosion resistance of the Al-Cu-Mg alloy. In addition, Al3Sc displayed good electrochemical compatibility with Al-alloys and should be expected to pose no greater a corrosion risk in Al alloys[32,39]. Norman et al.[40] found that the refinement effect in Al-Sc alloys occurred only with hypereutectic compositions (obove 0.55%). In our work, however, the size of sub-grains decreased 10
with increasing Sc and Zr alloying element additions in the Al-Cu-Mg alloy, as shown in Fig. 4. This indicated that the refinement effect manifested itself at lower Sc concentrations in the presence of Zr, which was consistent with the finding reported by Davydov et al.[17]. Furthermore, grain refinement can lead to the reduction of planar slip and promote a more homogeneous slip mode, which can effectively reduce inter-crystalline fracture and corrosion susceptibility. 4.2 Cross section morphology of inter-granular corrosion It was reported that there were two different corrosion processes associated with S (S’, S’’) phase with prolonged exposure to a corrosive environment[41], which is consistent with the corrosion behavior of Al2CuLi[42]. At the first stage of corrosion process, the potential of S (S’, S’’) phase is negative to the Al matrix. Therefore, the anodic dissolution occurred at S (S’, S’’) phase[43], forming a cross section morphology of inter-granular corrosion as shown in Fig.6 (j). With prolonging immersion time, active Mg in the S particles is preferentially dissolved, resulting in Cu-rich remnants[38]. Like Mg in S phase, Mg in Mg2Si particle also is preferentially dissolved[44]. With the dissolution of Mg element, the potential of S (S’, S’’) phase moves towards positive with respect to the Al matrix, resulting in the anodic dissolution of the Al base at the adjacent periphery of S (S’) phase[43], and then forming the cross section morphology of inter-granular corrosion as shown in Fig. 6(i). That is to say, there exists a corrosion mechanism conversion of the precipitate of S[41]. It can be inferred that the Cu/Mg ratio reduced after Sc and Zr alloying elements additions as shown in Table 4. The Cu/Mg ratio in Al-Cu-Mg alloy gradually increased from 10.16 to 30.54 with prolonging immersion time from 3 to 24 h in corrosive environment. However, the corresponding Cu/Mg ratio of Al-Cu-Mg-0.2Sc-0.4Zr alloy only augment from 2.07 to 6.56, indicating that adding minor Sc and Zr can retard the dissolution of active Mg during immersion in corrosive environment, which may delay the conversion of the corrosion mechanism of S phase. In addition, Cu content of the grain boundary precipitates decreased after adding Sc and Zr elements, evident from Table 2, is expected to have strong influence on inter-granular corrosion. The additions of Sc and Zr can restrict the dissolution of Mg and retard the segregation of Cu element, resulting in hindering the corrosion mechanism conversion of S phase, which finally presents the different cross section morphology between Al-Cu-Mg and Al-Cu-Mg-Sc-Zr alloy after inter-granular corrosion test. As the discussion in section 4.1, the discontinuous precipitate distribution at the sub-grain boundaries in Al-Cu-Mg-Sc-Zr alloy cut off the path of anode dissolution and decreased the 11
corrosion speed, which was another key factor in changing cross section morphology of Al-Cu-Mg-Sc-Zr alloy. 5. Conclusion Compared with the microstructures of peak-aged Al-Cu-Mg alloy, the degree of recrystallization was significantly decreased, the size of sub-grains was remarkably reduced, the content of Cu element decreased along the grain boundary and the spacing of grain boundary precipitates was substantially increased after adding minor Sc and Zr with a Sc/Zr ratio of 1/2. The additions of Sc and Zr to Al-Cu-Mg alloy changed cross section morphology of inter-granular corrosion from an undercutting to an elliptical shape due to the hindrance of the corrosion mechanism conversion of the S phase. Furthermore, corrosion resistance of Al-Cu-Mg alloy was enhanced by adding a trace amount of Sc and Zr. Optimized corrosion performance of Al-4Cu-2Mg (wt%) alloy was achieved by adding 0.2Sc and 0.4Zr. The superior corrosion resistant property from micro-alloying additions of Sc and Zr resulted from refined grains and the discontinuous distribution of grain boundary precipitates, which inhibited the recrystallization. The change in cross section morphology of the alloy with adding Sc and Zr was mainly attributed to the difference in electrochemical property of grain boundary precipitates, which prevented the anodic dissolution of the Al matrix. Acknowledgment This work was financially supported by the National High-tech Research & Development Program of China (2013AA031002).
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References [1] H.Z. Li, Y.J. Ou, H.J. Liao, X.P. Liang, J. Jiang, Mater. Sci. Forum 788 (2014) 208–214. [2] J.C. Williams, J.E.A Starke, Acta Mater. 51 (2003) 5775–5799. [3] A. Heinz, A. Haszler, C. Keidel, S. Moldenhauer, R. Benedictus, W.S. Miller, Mater. Sci. Eng. A 280 (2000) 102–107. [4] Y. Deng, Z.M. Yin, K. Zhao, J.Q. Duan, J. Hu, Z.B. He, Corros. Sci. 65 (2012) 288–298. [5] J. Liu, P. Yao, N.Q. Zhao, C.S. Shi, H.J. Li, X. Li, D.S. Xi, S. Yang, J. Alloys Compd. 657 (2016) 717–725 [6] Y.J. Shi, Q.L. Pan, M.J. Li, X. Huang, B. Li, J. Alloys Compd. 612 (2014) 42–50. [7] N.L. Sukiman, R.K. Gupta, R.G. Buchheit, N. Birbilis, Corros. Eng. Sci. Technol. 49 (2014) 254– 262. [8] T. Wang, T. Gao, P. Zhang, J.F. Nie, X.F. Liu, J. Alloys Compd. 589 (2014) 19–24. [9] R.K. Gupta, R. Zhang, C.H.J. Davies, N. Birbilis, Corrosion 70 (2014) 402–413. [10] R.K. Gupta, D. Fabijanic, T. Dorin, Y. Qiu, J.T. Wang, N. Birbilis, Mater. Des. 84 (2015) 270– 276. [11] Y. Wang, R.K. Gupta, N.L. Sukiman, R. Zhang. C.H.J. Davies, N. Birbilis, Corros. Sci. 73 (2013) 181–187. [12] N.Q. Tuan, A.C. Alves, F. Topan, A.B. Lopes, A.M.P. Pinto, Mater. Corros. 67 (2016) 60–71. [13] R.K. Gupta, Y. Wang, R. Fang, N.L. Sukiman, C.H.J Davies, N. Birbilis, Corrosion 69 (2013) 4– 8. [14] N.L. Sukiman, R.K. Gupta, R.G. Buchheit, N. Birbilis, Corros. Eng. Sci. Technol. 49 (2014) 263–268. [15] H.C. Fang, H. Chao, K.H. Chen, Effect of Zr, Mater. Sci. Eng. A 610 (2014) 10–16. [16] G.S. Peng, K.H. Chen, H.C. Fang, S.Y. Chen, Mater. Des. 36 (2012) 279–283. [17] V.G. Davydov, T.D. Rostova, V.V. Zakharov, Yu.A. Filatov, V.I. Yelagin, Mater. Sci. Eng. A 280 (2000) 30–36. [18] Y.C. Tzeng, C.T. Wu, S.L. Lee, Mater. Lett. 161 (2015) 340–342. [19] J.W. Zhao, C. Xu, G.Z. Dai, S.S. Wu, J. Han, Mater. Lett. 173 (2016) 22–25. [20] Y. Qian, J.L. Xue, Z.J. Wang, Z.H. Yang, P. Qian, JOM 68 (2016) 1293–1300. [21] G. Li, N.Q. Zhao, T. Liu, J.J. Li, C.N. He, C.S. Shi, E.Z. Liu, J.W. Sha, Mater. Sci. Eng. A 617 13
(2014) 219–227. [22] C. Xu, W.L Xiao, R.X. Zheng, S. Hanada, H. Yamagata, C.L. Ma, Mater. Des. 88 (2015) 485– 492. [23] H.F. Huang, F. Jiang, J. Zhou, L.L. Wei, J.P. Qu, L.L Liu, J. Mater. Eng. Perfor. 24 (2015) 4244– 4252. [24] B. Li, Q.L. Pan, C.P. Chen, H.H. Wu, Z.M. Yin, J. Alloys Compd. 664 (2016) 553–564. [25] A.M. Samuel, S.A. Alkahtani, H.W. Doty, F.H. Samuel, Mater. Des. 88 (2015) 1134–1144. [26] B Li, Q.L Pan, X. Huang, Z.M. Yin, Mater. Sci. Eng. A 616 (2014) 219–228. [27] L.Q. Zhu, A. Gu, H.C. Liu, J.Z. Liu, X.P. Ye, B.R. Hu, J. Aeronaut. Mater. 28 (2006) 61–66. [28] GB/T 7998-2005, National Standard of China, Test Method for Inter-granular Corrosion of Aluminum Alloys. [29] Y.C. Liu, Y.C. Xia, Y.Q. Jiang, L.T. Li, Mater. Sci. Eng. A 556 (2012) 796–800. [30] C. Li, Q.L. Pan, Y.J. Shi, Y. Wang, B. Li, Mater. Des. 55 (2014) 551–559. [31] L.S. Toropova, D.G. Eskin, M.L. Kharakterova, T.V. Dobatkina, Advanced Aluminium Alloys Containing Scandium—Structure and Properties Gordon and Breach Science Publisher, The Netherlands, 1998. [32] V. Neubert, B. Smola, I. Stulıkova, A. Bakkar, J. Reuter, Mater. Sci. Eng. A 464 (2007) 358– 364. [33] M. Vlach, I. Stulikova, B. Smola, T. Kekule, H. Kudrnova, V. Kodetova, V. Ocenasek, J. Malek, V. Neubert : Kovove Mater. 53 (2015) 295–304 . [34] T. Dorin, N. Stanford, N. Birbilis, R.K. Gupta, Corros. Sci. 100 (2015) 396–403. [35] R. Zhang, R.K. Gupta, C.H.J. Davies, A.M. Hodge, M. Tort, K. Xia, N. Birbilis, Corrosion 72 (2016) 160–168. [36] R.K. Gupta, N.L. Sukiman, M.K. Cavanaugh, B.R.W. Hinton, C.R. Hutchinson, N. Birbilis, Electro. Acta 66 (2012) 245–254. [37] J.A. Lyndon, R.K. Gupta, M.A. Gibson, N.Birbilis, Corros. Sci. 70 (2013) 290–293. [38] J.F. Li, Z.Q. Zheng, N. Jiang, C.Y. Yan, Mater. Chem. Phys. 91 (2005) 325–329. [39] M.K. Cavanaugh, N. Birbilis, R.G. Buchheita, F. Bovard, Scr. Mater. 56 (2007) 995–998. [40] A .F. Norman, P.B. Prangnell, R.S. Mcewen, Acta Mater. 46 (1998) 5715–5732. [41] J.C. Li, N. Birbilis, R.G. Buchheit, Corros. Sci. 101 (2015) 155–164. 14
[42] J.F. Li, Z.Q. Zheng, S.C. Li, W.J. Chen, W.D. Ren, X.S.Zhen, Corros. Sci. 49 (2007) 2436– 2449. [43] M.H. Shao, Y. Fu, R.G. Hu, C.J. Lin, Mater. Sci. Eng. A 344 (2003) 323–327. [44] K. Mizuno, A. Nylund, I. Olefjord, Corros. Sci. 43 (2001) 381–396.
15
Figure captions Fig. 1 Hardness of the Al-Cu-Mg-Sc-Zr alloys aged at 190 oC. Fig. 2 Optical microstructures of the alloy in solid–solution state: (a) Al-Cu-Mg alloy; (b) Al-Cu-Mg-0.1Sc-0.2Zr alloy; (c) Al-Cu-Mg-0.2Sc-0.4Zr alloy. Fig.3 SEM-EBSD measurements of alloys peak-aged at 190oC. Orientation maps: (a) Al-Cu-Mg; (b) Al-Cu-Mg-0.1Sc-0.2Zr; (c) Al-Cu-Mg-0.2Sc-0.4Zr; Misorientation angle distribution: (d) Al-Cu-Mg, (e) Al-Cu-Mg-0.1Sc-0.2Zr, (f) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.4 TEM images of the studied alloys with different amount of Sc and Zr additions: (a) Al-Cu-Mg alloy; (b) Al-Cu-Mg-0.1Sc-0.2Zr alloy; (c) Al-Cu-Mg-0.2Sc-0.4Zr alloy; (d and e) Al3(Sc, Zr) in Al-Cu-Mg-0.2Sc-0.4Zr alloy; (f) SAD in [001]Al projection; (g) SAD in [011]Al projection. Fig.5 Metallographic images of cross-section morphology after immersion in the corrosive environment: (a) Al-Cu-Mg; (b) Al-Cu-Mg-0.1Sc-0.2Zr; (c) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.6 Metallographic images of cross-section morphology of the peak-aged alloys show the developed corrosion pits morphology after exposure to IGC environment for different durations: Al-Cu-Mg alloy exposed for (a) 3 h; (c) 6 h; (e) 12 h and (g) 24 h; Al-Cu-Mg-0.2Sc-0.4Zr alloy exposed for (b) 3 h; (d) 6 h; (f) 12 h and (h) 24 h; Schematic drawing of pit shape of (i)Al-Cu-Mg alloy; (j) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.7 SEM images of (a) overall topography of the cross-section morphology; (b) high magnification of the cross-section morphology.
16
Figure List
Fig. 1
Fig. 2
Fig. 3
17
Fig. 4
18
Fig. 5
Fig. 6
Fig. 7
19
Table 1 Chemical compositions of studied alloys (in wt%), as determined by inductively coupled plasma–optical emission spectrometry (ICP–OES) Alloys
Cu
Mg
Zr
Sc
Fe
Si
Sc+Zr
Al
Al-Cu-Mg
3.92
2.11
0
0
0.17
0.06
0
Bal.
Al-Cu-Mg-0.1Sc-0.2Zr
4.12
1.89
0.23
0.09
0.16
0.07
0.32
Bal.
Al-Cu-Mg-0.2Sc-0.4Zr
3.96
1.86
0.41
0.18
0.17
0.07
0.59
Bal.
Table 2 Average EDS test results (three tests) of the grain boundary precipitates in Al-Cu-Mg-Sc-Zr alloys (wt%)
Element Al-Cu-Mg Al-Cu-Mg-0.1Sc-0.2Zr Al-Cu-Mg-0.2Sc-0.4Zr Cu
21.79
11.88
10.54
Mg
8.72
7.45
6.97
Cu:Mg
2.50
1.59
1.51
Table 3 Variations of the σb and δb with different amount of Sc and Zr addition for the peak-aged Al-Cu-Mg alloys
Alloys
Al-Cu-Mg
Al-Cu-Mg-0.1Sc-0.2Zr
Al-Cu-Mg-0.2Sc-0.4Zr
σb
σf
δUTS
δb
δf
δEI
(MPa)
(MPa)
(%)
(%)
(%)
(%)
363.5
320.3
11.88
7.36
2.64
64.13
(±17)
(±21)
(±0.1)
(±0.2)
436.0
411.5
13.64
10.32
(±14)
(±19)
(±0.2)
(±0.1)
432.5
412.6
14.16
12.36
(±18)
(±10)
(±0.1)
(±0.3)
20
5.64
4.60
24.34
12.71
Table 4 Average EDS test results (three tests) at the bottom of inter-granular corrosion of Al-Cu-Mg-Sc-Zr alloys after immersion for different time (wt%) Element
Al-Cu-Mg
Al-Cu-Mg-0.2Sc-0.4Zr
3h
3h
24 h
24 h
Cu
29.13 51.30
10.65
32.49
Mg
2.86
5.14
4.95
Cu:Mg
10.16 30.54
2.07
6.56
1.68
21
S1005-0302(17)30004-X http://dx.doi.org/doi:10.1016/j.jmst.2016.12.003 JMST 883
To appear in: Received date: Revised date: Accepted date:
6-9-2016 31-10-2016 19-12-2016
Please cite this article as: {http://dx.doi.org/
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Effect of Sc and Zr Additions on Microstructures and Corrosion Behavior of Al-Cu-Mg-Sc-Zr Alloys Fangfang Sun1, Guiru Liu Nash2, Qunying Li1, Enzuo Liu1,3, Chunnain He1,3, Chunsheng Shi1,3, Naiqin Zhao1,3* 1
School of Materials Science and Engineering, Tianjin University, Tianjin 300072, China
2
Department of Mechanical, Materials and Aerospace Engineering, Illinois Institute of Technology,
USA 3
Tianjin Key Laboratory of Composite and Functional Materials, Tianjin 300072, China
*Corresponding author. Tel. & Fax: +86 1 27891371. E-mail address: [email protected] (N. Zhao).
[Received 6 September 2016; Received in revised form 31 October 2016; Accepted 19 December 2016]
Graphical Abstract
1
Highlights: 1.
Adding Sc and Zr changed the cross section morphology of inter-granular corrosion.
2.
There exists a corrosion mechanism conversion of S precipitates in Al-Cu-Mg alloy.
3.
Adding Sc and Zr restricted the corrosion mechanism conversion of S particles.
4.
The Al-Cu-Mg-Sc-Zr alloy possesses excellent inter-granular corrosion properties.
5.
The content of Cu and Mg elements decreased along the grain boundary after minor Sc and Zr addition.
2
The effects of adding the alloy element Sc to Al alloys on strengthening, recrystallization and modification of the grain microstructure have been investigated. The combination of Sc and Zr alloying not only produces a remarkable synergistic effect of inhibition of recrystallization and refinement of grain size but also substantially reduce the amount of high-cost additional Sc. In this work, the microstructures and corrosion behavior of a new type of Al-Cu-Mg-Sc-Zr alloy with Sc/Zr ratio of 1/2 were investigated. The experimental results showed that the Sc and Zr additions to Al-Cu-Mg alloy could strongly inhibit recrystallization, refine grain size, impede the segregation of Cu element along the grain boundary and increase the spacing of grain boundary precipitates. In addition, adding Sc and Zr to Al-Cu-Mg alloy effectively restricts the corrosion mechanism conversion associated with Al2CuMg particles, which resulted in the change of the cross-section morphology of inter-granular corrosion from an undercutting to an elliptical shape. The susceptibility to inter-granular corrosion was significantly decreased with increasing Sc and Zr additions to the Al-Cu-Mg alloy. The relationships between microstructures evolution and inter-granular corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys were also discussed.
Key words: Aluminum alloy; Microstructure; Scandium; Inter-granular corrosion
1. Introduction Al-Cu-Mg alloys are important structural materials that have been widely used in the aerospace field due to the high specific strength, excellent heat resistance and easy processability[1-3]. However, the tendency for inter-granular corrosion (IGC) of Al-Cu-Mg alloys is relatively high due to the potential difference between the heterogeneous distributions of S (Al2CuMg) or θ (Al2Cu) precipitates along the grain boundaries and in the matrix. Therefore, development of high specific strength with excellent corrosion resistance of Al-Cu-Mg alloys is a main research goal for potential applications in aviation, aerospace, and other high-technology sectors[4]. Micro-alloying with rare earth or transition metals (such as Sc, Zr, Cr, Nb, Sr and Ti) was discovered as an important method to modify microstructures and improve performance of the aluminum alloy[5-11]. Compared to the Al-4Mg alloy, the addition of Sc and Yb reduced the grain size, 3
changed the morphology of intermetallic particles and decreased the corrosion potential values of Al-4Mg alloy in 3.5 wt% NaCl solution[12]. Gupta et al.[13,14] studied the sensitization behavior of Al-5Mg alloy using nitric acid mass loss test found that micro-alloying additions Nd ( >0.11%), Sr, Si, and Ti significantly reduced IGC susceptibility of Al-5Mg alloy. Moreover, Fang et al.[15] and Peng et al.[16] showed that complex additions of Er, Zr and Cr to Al-Zn-Mg-Cu alloys remarkably enhanced resistance to IGC, which was contributed by the formation of 20–50 nm particles strongly inhibiting the recrystallization behavior of the Al matrix. However, Sc was reported to be the most effective micro-alloying element to modify microstructures and improve properties of Al alloy[17]. The quench sensitivity of the Al-7Si-0.6 Mg alloy was effectively improved by more than 60% with 0.04% Sc additions[18]. The effects of Sc additions on microstructure and properties of rheo-diecasting wrought aluminum alloy were investigated by Zhao et al.[19]. This study concluded that Sc additions not only influenced the morphology of the primary α-Al particles, but also promoted the formation of fine α-Al globules during the solidification process of the remaining liquid, resulting in the considerable improvement in mechanical properties. Moreover, Sc-containing Al alloys are limited by high cost. Previous studies had demonstrated that the combined alloying elements of Sc and Zr not only produced remarkable synergistic effects but also substantially reduced the Sc content. Qian et al.[20] systematically investigated the mechanical properties of L12-Al3(Sc,Zr) using first principles calculations. The results showed that the ideal strength and ductility of L12- Al3(Sc,Zr) increased with increasing Zr addition. Li et al.[21] investigated different Sc/Zr ratio addition to the Al-Zn-Mg-Cu alloy and concluded that the strengthening effect was better with a Sc/Zr ratio of 2/4 than that without Sc and Zr additions or only adding Sc. Most of the literature has focused on microstructure and mechanical properties, while studies on corrosion behavior of Sc- and Zr-containing Al alloys, especially of Sc- and Zr-containing Al-Cu-Mg alloys, are scarce[5,12,21-25]. The improvements of mechanical properties of Al-7Si-0.65Mg alloy were attributed to microstructural refinement, particularly the modification of eutectic Si and precipitation of secondary nano-scale Al3(Sc, Zr) dispersoids after minor Sc and Zr additions[22]. The effect of Sc and Zr on corrosion behavior of Al-Mg alloys[9, 12] and Al-Zn-Mg alloys[4, 6, 15, 26] have been reported. Li et al.[26] pointed out that the enhancement of corrosion performance of Al-Zn-Mg-Mn alloy was due to the effect of inhibiting recrystallization behavior caused by Al3 (Sc, Zr) precipitates and interrupted distributions of grain boundary precipitates. Little research has been made on effects of 4
Sc and Zr alloying element additions on microstructure and corrosion properties of Al-Cu-Mg alloys. Furthermore, Zhu et al.[27] deemed that AA2024 (Al-Cu-Mg alloys) displayed a completely different morphology of corrosion cross-section from 7B04 aluminum alloys. The relationship between the corrosion property, microstructure and the cross-section morphology of the inter-granular corrosion with Sc and Zr alloying element additions to Al-Cu-Mg alloys is still unclear. Therefore, it is necessary to study the relationship between corrosion resistance (including cross-section morphology of inter-granular corrosion) and micro-alloying element addition to understand the corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys. The purpose of this work is to conduct a comprehensive study the effect of a specific Sc/Zr ratio on microstructure and corrosion behavior of Al-Cu-Mg alloys. The microstructure evolution, cross-section morphology of inter-granular corrosion and corrosion mechanism of Al-Cu-Mg-Sc-Zr alloys will be discussed in this paper. 2. Materials and Experimental Methods 2.1 Materials preparation Three types of Al-Cu-Mg (in wt%) alloys with Sc and Zr additions at a constant Sc/Zr ratio of 1/2 were prepared. In order to fully nucleate Al3(Sc, Zr) particles, we designed the alloy with a slight excess Zr content, since the Al3Sc phase nucleates on the surface of the Al3Zr phase particle. Flat bar ingots with 20 mm in thickness, were provided by Hunan Rare-Earth Metal Research Institute, China. The chemical compositions of the alloys were analyzed by inductively coupled plasma–optical emission spectrometry as shown in Table 1. The ingots were homogenized at 480 oC for 20 h, hot rolled from 20 to 6.9 mm and then cold rolled to sheets with a 2.7 mm final thickness. The total deformation reduction of each alloy was 86.5%. The sheets of all three alloys were solution treated at 500 oC for 1 h, subsequently water quenched and artificially aged at 190 oC for various time. Ten hardness measurements were conducted on each sample and the average results with the standard deviation from the mean were reported. Aging hardening curves of the Al-Cu-Mg(-Sc-Zr) alloys at 190 oC are presented in Fig. 1. It was notable that all of the standard deviations of the samples were below HV3.5. The hardening behavior of the alloys aged at 190 oC displayed a dual-peak phenomenon. After minor Sc and Zr addition, the second peak-aged time was reduced and the peak hardness was increased with increasing amounts of Sc and Zr. The peak-ageing time of Al-Cu-Mg, Al-Cu-Mg-0.3(Sc+Zr) and Al-Cu-Mg-0.6(Sc+Zr) alloys were 22, 10, and 10 h, respectively. 5
Hereafter, all of the alloys are in the peak-aged condition unless otherwise specified. 2.2 Microstructural characterization Metallographic observation was conducted for the solid–solution state samples using a Leica optical microscope. The samples were prepared by conventional mechanical polishing and then etched with Keller's reagent (95 vol.% H2O + 2.5 vol.% HNO3 + 1.5 vol.% HCl + 1 vol.% HF). Misorientation angles between grains/sub-grains of T6 treated specimens were measured by electron back scattered diffraction (EBSD). The microstructure of peak-aged sheets was analyzed by transmission electron microscopy (TEM). Energy dispersive spectroscopy (EDS) was utilized to investigate the composition of grain boundary precipitates. Thin foils for TEM observations were prepared by twin-jet electro-polishing at 42.5 mA in a solution of 30% nitric acid and 70% methanol solution cooled to -30 oC. A JEM-2100F electron microscope and a TECNAI G2 F20 electron microscope were used to conduct the analysis. 2.3 Corrosion tests After grinding and polishing, all the peak-aged specimens were immersed in the corrosive solution (57 g NaCl + 1 l H2O + 10 ml H2O2) at (35±2) oC for 6 h. The ratio of the exposed area of specimens and the volume of corrosion solution was less than 20 mm2/ml, according to inter-granular corrosion testing requirement of GB/T 7998-2005[28]. In order to evaluate the mechanical properties, tensile tests were conducted using the samples with and without immersing in the corrosive environment. Tensile tests were conducted using a CSS-44100 electronic universal testing machine with 1 mm/min loading speed. The axis of the tensile specimens was parallel to the rolling direction of the thin sheets. All experiments were conducted three times and the average testing results are reported with the standard deviation. The cross-section (perpendicular to the rolling direction) of samples was measured using a Leica optical microscope after the inter-granular corrosion tests. After removing the corrosion products from the test specimens, energy dispersive spectroscopy (SEM-EDS) was utilized to investigate the composition of the region at the bottom of inter-granular corrosion with prolonged time for immersion in the corrosive solution. 3. Results 3.1
Microstructures Fig. 2 shows the microstructures of three different alloys after solution treatment at 500 oC for 1
h. It was found that the grain size was reduced with increasing amounts of Sc and Zr in the 6
Al-Cu-Mg alloy. However, it showed that Al-Cu-Mg-0.1Sc-0.2Zr alloy had slightly smaller banded structure than Al-Cu-Mg-0.2Sc-0.4Zr alloy. Fig. 3 shows the SEM-EBSD images and misorientation angle distribution of three alloys. It was found that the Al-Cu-Mg alloy had uniaxial grains and the alloys with additional Sc and Zr had elongated grains as shown in Fig. 3(a, b and c). The boundary misorientation angle distributions of Al-Cu-Mg and Al-Cu-Mg-Sc-Zr alloys after peak ageing (190 o
C) are shown in Fig. 3(d, e and f). After the solution treatment at 500 oC for 1 h, full
recrystallization occurred in the Al-Cu-Mg alloy resulting in high angle grain boundaries (HAGBs) with the misorientation angles mainly ranging from 20° to 55° (Fig. 3(a and d)). However, Al-Cu-Mg alloy with Sc and Zr alloying element additions still had a fibrous unrecrystallized microstructure. The grain boundaries were characterized as low angle grain boundaries (LAGBs) with the misorientation angles mostly less than 10°. In addition, with the increase in the Sc and Zr alloying element additions, the proportion of LAGBs increased, indicating that the effect of inhibiting recrystallization was improved (Fig. 3(e and f)). The above observation was further verified by TEM images as shown in Fig. 4. It was found that the Al-Cu-Mg-Sc-Zr peak-aged alloys contained more micro-scale sub-grains than the recrystallized Al-Cu-Mg alloy as shown in Fig. 4(a, b and c). The grain boundary precipitates (GBPs) in Al-Cu-Mg alloy were distributed continuously, as shown in Fig. 4(a). The GBPs were distributed discontinuously by minor Sc and Zr additions as shown in the upper right corner of Fig. 4(b and c). In addition, the interparticle spacing of GBPs was increased with the increase in the Sc and Zr additions, as shown in Fig. 4(a, b and c). It was also found that the size of sub-grains decreased with the increase in the Sc and Zr alloying element additions as shown in Fig. 4(b and c). Very fine precipitate particles were found inside the sub-grains as shown in Fig. 4(d and e). The [001] selected area diffraction (SAD) pattern revealed that the fine precipitation particles were Al3(Sc,Zr) as shown in Fig. 4(f). The sub-grains might have resulted from the formation of nanometer sized, coherent, secondary Al3(Sc,Zr) particles with the additions of Sc and Zr. The [011] Al SAD pattern (Fig. 4(g)) revealed that light spots located in 1/3<200> and 2/3<200> directions corresponding to acicular exudation phase was S’ phase[29]. Thus, S’ phase is the main strengthening precipitate in the peak-aged condition of Al-Cu-Mg alloy. The chemical compositions of GBPs determined by TEM-EDS are listed in Table 2. It was found that the percentage of the Cu and Mg elements in GBPs were substantially decreased with the 7
additions of Sc and Zr, which implied that the segregation of Cu and Mg elements along the grain boundary was inhibited to some extent by minor Sc and Zr additions. This finding was consistent with previous results reported by Wang et al.[8], which supported the finding that the grain refinement was an effective method to reduce or eliminate grain boundary segregation. Furthermore, the value of the Cu/Mg ratio in the precipitates significantly decreased with increasing amount of Sc and Zr additions. 3.2 Inter-granular corrosion The cross section morphologies of inter-granular corrosion were analyzed by optical microscopy, as shown in Fig. 5. It was found in Fig. 5 that the maximum depth of inter-granular corrosion was 124 m for the Al-Cu-Mg alloy, 107 m for the Al-Cu-Mg-0.1Sc-0.2Zr alloy and 61m for the Al-Cu-Mg-0.2Sc-0.4Zr alloy. The dramatic decrease in the maximum depth of inter-granular corrosion indicated that the inter-granular corrosion resistance was greatly increased with minor Sc and Zr additions. In order to evaluate the susceptibility to inter-granular corrosion of the studied alloys, the ultimate tensile strength loss ratio (δUTS) and elongation loss ratio (δEI) were calculated as:
UTS
b - f b
(1)
EI
b - f b
(2)
where σb and δb are the ultimate tensile strength and total elongation of samples without immersion in the corrosive solution, respectively; σf and δf are the ultimate tensile strength and total elongation of the samples after immersion in the corrosive solution for 6 h at 35 oC. The variations of the δUTS and δEI with different amounts of Sc and Zr additions for the peak-aged Al-Cu-Mg alloys are listed in Table 3. It was found that the δUTS and δEI decreased with increasing amount of Sc and Zr additions, which indicated that the minor additions of Sc and Zr tended to decrease the inter-granular corrosion susceptibility of the Al-Cu-Mg alloys. The observation was consistent with the measured maximum depth of inter-granular corrosion as a function of the amount of Sc and Zr alloying element additions. Furthermore, the decrease in ultimate tensile strength of Al-Cu-Mg-0.2Sc-0.4Zr alloy might have resulted from the formation of coarse and primary Al3Zr particles when the content of Zr was over 0.3 wt%[30]. 8
3.3 Evolution of the cross section morphology of inter-granular corrosion The cross section morphology of the tested samples contains abundant characteristic information including corrosion types, corrosion mechanism, microstructure and residual fatigue strength[27]. A comprehensive investigation of the effect of minor Sc and Zr additions on the evolution of cross section morphology of inter-granular corrosion and corrosion mechanism has been carried out. Fig. 6 shows the cross section morphology of the Al-Cu-Mg and Al-Cu-Mg-0.2Sc-0.4Zr alloys after soaking in the inter-granular corrosion solution for different durations. Compared to Al-Cu-Mg alloy, it was found that corrosion area and corrosion depth for the same soaking duration were both significantly reduced by adding minor Sc and Zr elements. The cross section morphology of inter-granular corrosion of the Al-Cu-Mg alloy appeared to extend in a lateral direction below the surface. Therefore, the surface width of the cross section morphology in the Al-Cu-Mg alloy was small with a bigger inner width beneath the surface as shown in Fig. 6(a, c, e and g), which would have resulted in the miscalculation of the corrosion degree. However, the surface width of the cross-section morphology of inter-granular corrosion in the Al-Cu-Mg-0.2Sc-0.4Zr alloy was the largest and gradually reduced toward the bottom as shown in Fig. 6(b, d, f and h). A schematic of the cross section morphology of the two alloys are illustrated in Fig. 6(i and j). The overall topography of the cross section morphology is shown in Fig. 7. It was found that the edge of the cross section of inter-granular corrosion was smooth and bottom of cross section was uneven. Table 4 shows the chemical compositions at the bottoms of cross section of the Al-Cu-Mg and Al-Cu-Mg-0.2Sc-0.4Zr alloys, respectively. It was found that Cu was segregated at the bottom of the cross section in the Al-Cu-Mg alloy after a prolonged immersion in the corrosive solution. Compared to the Al-Cu-Mg alloy, the Cu and the Cu/Mg ratio were both significantly reduced at the bottom of cross section in the Al-Cu-Mg-0.2Sc-0.4Zr alloy for the same soaking time. 4. Discussion 4.1 Corrosion mechanism The combination of Sc and Zr significantly enhanced the resistance to recrystallization of the Al-Cu-Mg alloy, which can be attributed to the formation of high density, coherent and dispersed secondary Al3(Sc, Zr) particles. The Al3(Sc,Zr) particles were identified by both the presence of Ashby-Brown contrast (Fig. 4(d and e)) and the superstructure reflections like [001]Al (Fig. 4(f))[15]. 9
The coherent and fine dispersed secondary Al3(Sc,Zr) particles strongly pin dislocations and the sub-grain/grain boundaries during further heat treatment processes, resulting in unrecrystallized sub-grain structures retained in the deformed Al and Al alloys[5,21,31,32]. Vlach et al.[33] reported that combination of Sc and Zr additions to Al suppressed recrystallization during annealing at 550 oC. In our work, the effect of recrystallization inhibition was increased with increasing amount of Sc and Zr additions. It is well documented that the major microstructural features affecting the inter-granular corrosion resistance are the morphology and chemical composition of the GBPs[26, 34-37]. With the addition of Sc and Zr, the fiber-like sub-grains consisting of LAGBs were obtained. Compared to LAGBs, HAGBs have high energy, which is apt to form continuous and coarse precipitates along the grain boundaries. However, the LAGBs have low energy, which is prone to the formation of finer and discontinuous GBPs along the low angle unrecrystallized sub-grain boundaries. According to the electrochemical behavior of the precipitates in Al-Cu-Mg alloy, the potentials of S (S’, S’’) phase is more negative than the matrix and precipitate free zone (PFZ)[38]. Thus, S (S’, S’’) phase would act as the anode and firstly dissolve when the alloy is exposed to a corrosive environment according to the anodic dissolution mechanism. Therefore, the continuous S’ phase distribution in fully recrystallized Al-Cu-Mg alloy offers an unblocked path for corrosion. However, the discontinuous distribution of S’ precipitates along LAGBs could cut off the path of anode dissolution in the Al-Cu-Mg-Sc-Zr alloys, which results in the decrease in the corrosion speed and the improvement of corrosion resistance. In summary, adding Sc and Zr alloying elements to Al-Cu-Mg alloys can enhance the resistance to recrystallization and stabilize the deformed fibrous microstructure with numerous LAGBs due to the formation of Al3(Sc,Zr) particles. The deformed fibrous microstructure with the discontinuous precipitate distribution at the sub-grain boundaries resulted in the improvement in corrosion resistance. The more the Sc and Zr additions, the more remarkable the effect on improving corrosion resistance of the Al-Cu-Mg alloy. In addition, Al3Sc displayed good electrochemical compatibility with Al-alloys and should be expected to pose no greater a corrosion risk in Al alloys[32,39]. Norman et al.[40] found that the refinement effect in Al-Sc alloys occurred only with hypereutectic compositions (obove 0.55%). In our work, however, the size of sub-grains decreased 10
with increasing Sc and Zr alloying element additions in the Al-Cu-Mg alloy, as shown in Fig. 4. This indicated that the refinement effect manifested itself at lower Sc concentrations in the presence of Zr, which was consistent with the finding reported by Davydov et al.[17]. Furthermore, grain refinement can lead to the reduction of planar slip and promote a more homogeneous slip mode, which can effectively reduce inter-crystalline fracture and corrosion susceptibility. 4.2 Cross section morphology of inter-granular corrosion It was reported that there were two different corrosion processes associated with S (S’, S’’) phase with prolonged exposure to a corrosive environment[41], which is consistent with the corrosion behavior of Al2CuLi[42]. At the first stage of corrosion process, the potential of S (S’, S’’) phase is negative to the Al matrix. Therefore, the anodic dissolution occurred at S (S’, S’’) phase[43], forming a cross section morphology of inter-granular corrosion as shown in Fig.6 (j). With prolonging immersion time, active Mg in the S particles is preferentially dissolved, resulting in Cu-rich remnants[38]. Like Mg in S phase, Mg in Mg2Si particle also is preferentially dissolved[44]. With the dissolution of Mg element, the potential of S (S’, S’’) phase moves towards positive with respect to the Al matrix, resulting in the anodic dissolution of the Al base at the adjacent periphery of S (S’) phase[43], and then forming the cross section morphology of inter-granular corrosion as shown in Fig. 6(i). That is to say, there exists a corrosion mechanism conversion of the precipitate of S[41]. It can be inferred that the Cu/Mg ratio reduced after Sc and Zr alloying elements additions as shown in Table 4. The Cu/Mg ratio in Al-Cu-Mg alloy gradually increased from 10.16 to 30.54 with prolonging immersion time from 3 to 24 h in corrosive environment. However, the corresponding Cu/Mg ratio of Al-Cu-Mg-0.2Sc-0.4Zr alloy only augment from 2.07 to 6.56, indicating that adding minor Sc and Zr can retard the dissolution of active Mg during immersion in corrosive environment, which may delay the conversion of the corrosion mechanism of S phase. In addition, Cu content of the grain boundary precipitates decreased after adding Sc and Zr elements, evident from Table 2, is expected to have strong influence on inter-granular corrosion. The additions of Sc and Zr can restrict the dissolution of Mg and retard the segregation of Cu element, resulting in hindering the corrosion mechanism conversion of S phase, which finally presents the different cross section morphology between Al-Cu-Mg and Al-Cu-Mg-Sc-Zr alloy after inter-granular corrosion test. As the discussion in section 4.1, the discontinuous precipitate distribution at the sub-grain boundaries in Al-Cu-Mg-Sc-Zr alloy cut off the path of anode dissolution and decreased the 11
corrosion speed, which was another key factor in changing cross section morphology of Al-Cu-Mg-Sc-Zr alloy. 5. Conclusion Compared with the microstructures of peak-aged Al-Cu-Mg alloy, the degree of recrystallization was significantly decreased, the size of sub-grains was remarkably reduced, the content of Cu element decreased along the grain boundary and the spacing of grain boundary precipitates was substantially increased after adding minor Sc and Zr with a Sc/Zr ratio of 1/2. The additions of Sc and Zr to Al-Cu-Mg alloy changed cross section morphology of inter-granular corrosion from an undercutting to an elliptical shape due to the hindrance of the corrosion mechanism conversion of the S phase. Furthermore, corrosion resistance of Al-Cu-Mg alloy was enhanced by adding a trace amount of Sc and Zr. Optimized corrosion performance of Al-4Cu-2Mg (wt%) alloy was achieved by adding 0.2Sc and 0.4Zr. The superior corrosion resistant property from micro-alloying additions of Sc and Zr resulted from refined grains and the discontinuous distribution of grain boundary precipitates, which inhibited the recrystallization. The change in cross section morphology of the alloy with adding Sc and Zr was mainly attributed to the difference in electrochemical property of grain boundary precipitates, which prevented the anodic dissolution of the Al matrix. Acknowledgment This work was financially supported by the National High-tech Research & Development Program of China (2013AA031002).
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References [1] H.Z. Li, Y.J. Ou, H.J. Liao, X.P. Liang, J. Jiang, Mater. Sci. Forum 788 (2014) 208–214. [2] J.C. Williams, J.E.A Starke, Acta Mater. 51 (2003) 5775–5799. [3] A. Heinz, A. Haszler, C. Keidel, S. Moldenhauer, R. Benedictus, W.S. Miller, Mater. Sci. Eng. A 280 (2000) 102–107. [4] Y. Deng, Z.M. Yin, K. Zhao, J.Q. Duan, J. Hu, Z.B. He, Corros. Sci. 65 (2012) 288–298. [5] J. Liu, P. Yao, N.Q. Zhao, C.S. Shi, H.J. Li, X. Li, D.S. Xi, S. Yang, J. Alloys Compd. 657 (2016) 717–725 [6] Y.J. Shi, Q.L. Pan, M.J. Li, X. Huang, B. Li, J. Alloys Compd. 612 (2014) 42–50. [7] N.L. Sukiman, R.K. Gupta, R.G. Buchheit, N. Birbilis, Corros. Eng. Sci. Technol. 49 (2014) 254– 262. [8] T. Wang, T. Gao, P. Zhang, J.F. Nie, X.F. Liu, J. Alloys Compd. 589 (2014) 19–24. [9] R.K. Gupta, R. Zhang, C.H.J. Davies, N. Birbilis, Corrosion 70 (2014) 402–413. [10] R.K. Gupta, D. Fabijanic, T. Dorin, Y. Qiu, J.T. Wang, N. Birbilis, Mater. Des. 84 (2015) 270– 276. [11] Y. Wang, R.K. Gupta, N.L. Sukiman, R. Zhang. C.H.J. Davies, N. Birbilis, Corros. Sci. 73 (2013) 181–187. [12] N.Q. Tuan, A.C. Alves, F. Topan, A.B. Lopes, A.M.P. Pinto, Mater. Corros. 67 (2016) 60–71. [13] R.K. Gupta, Y. Wang, R. Fang, N.L. Sukiman, C.H.J Davies, N. Birbilis, Corrosion 69 (2013) 4– 8. [14] N.L. Sukiman, R.K. Gupta, R.G. Buchheit, N. Birbilis, Corros. Eng. Sci. Technol. 49 (2014) 263–268. [15] H.C. Fang, H. Chao, K.H. Chen, Effect of Zr, Mater. Sci. Eng. A 610 (2014) 10–16. [16] G.S. Peng, K.H. Chen, H.C. Fang, S.Y. Chen, Mater. Des. 36 (2012) 279–283. [17] V.G. Davydov, T.D. Rostova, V.V. Zakharov, Yu.A. Filatov, V.I. Yelagin, Mater. Sci. Eng. A 280 (2000) 30–36. [18] Y.C. Tzeng, C.T. Wu, S.L. Lee, Mater. Lett. 161 (2015) 340–342. [19] J.W. Zhao, C. Xu, G.Z. Dai, S.S. Wu, J. Han, Mater. Lett. 173 (2016) 22–25. [20] Y. Qian, J.L. Xue, Z.J. Wang, Z.H. Yang, P. Qian, JOM 68 (2016) 1293–1300. [21] G. Li, N.Q. Zhao, T. Liu, J.J. Li, C.N. He, C.S. Shi, E.Z. Liu, J.W. Sha, Mater. Sci. Eng. A 617 13
(2014) 219–227. [22] C. Xu, W.L Xiao, R.X. Zheng, S. Hanada, H. Yamagata, C.L. Ma, Mater. Des. 88 (2015) 485– 492. [23] H.F. Huang, F. Jiang, J. Zhou, L.L. Wei, J.P. Qu, L.L Liu, J. Mater. Eng. Perfor. 24 (2015) 4244– 4252. [24] B. Li, Q.L. Pan, C.P. Chen, H.H. Wu, Z.M. Yin, J. Alloys Compd. 664 (2016) 553–564. [25] A.M. Samuel, S.A. Alkahtani, H.W. Doty, F.H. Samuel, Mater. Des. 88 (2015) 1134–1144. [26] B Li, Q.L Pan, X. Huang, Z.M. Yin, Mater. Sci. Eng. A 616 (2014) 219–228. [27] L.Q. Zhu, A. Gu, H.C. Liu, J.Z. Liu, X.P. Ye, B.R. Hu, J. Aeronaut. Mater. 28 (2006) 61–66. [28] GB/T 7998-2005, National Standard of China, Test Method for Inter-granular Corrosion of Aluminum Alloys. [29] Y.C. Liu, Y.C. Xia, Y.Q. Jiang, L.T. Li, Mater. Sci. Eng. A 556 (2012) 796–800. [30] C. Li, Q.L. Pan, Y.J. Shi, Y. Wang, B. Li, Mater. Des. 55 (2014) 551–559. [31] L.S. Toropova, D.G. Eskin, M.L. Kharakterova, T.V. Dobatkina, Advanced Aluminium Alloys Containing Scandium—Structure and Properties Gordon and Breach Science Publisher, The Netherlands, 1998. [32] V. Neubert, B. Smola, I. Stulıkova, A. Bakkar, J. Reuter, Mater. Sci. Eng. A 464 (2007) 358– 364. [33] M. Vlach, I. Stulikova, B. Smola, T. Kekule, H. Kudrnova, V. Kodetova, V. Ocenasek, J. Malek, V. Neubert : Kovove Mater. 53 (2015) 295–304 . [34] T. Dorin, N. Stanford, N. Birbilis, R.K. Gupta, Corros. Sci. 100 (2015) 396–403. [35] R. Zhang, R.K. Gupta, C.H.J. Davies, A.M. Hodge, M. Tort, K. Xia, N. Birbilis, Corrosion 72 (2016) 160–168. [36] R.K. Gupta, N.L. Sukiman, M.K. Cavanaugh, B.R.W. Hinton, C.R. Hutchinson, N. Birbilis, Electro. Acta 66 (2012) 245–254. [37] J.A. Lyndon, R.K. Gupta, M.A. Gibson, N.Birbilis, Corros. Sci. 70 (2013) 290–293. [38] J.F. Li, Z.Q. Zheng, N. Jiang, C.Y. Yan, Mater. Chem. Phys. 91 (2005) 325–329. [39] M.K. Cavanaugh, N. Birbilis, R.G. Buchheita, F. Bovard, Scr. Mater. 56 (2007) 995–998. [40] A .F. Norman, P.B. Prangnell, R.S. Mcewen, Acta Mater. 46 (1998) 5715–5732. [41] J.C. Li, N. Birbilis, R.G. Buchheit, Corros. Sci. 101 (2015) 155–164. 14
[42] J.F. Li, Z.Q. Zheng, S.C. Li, W.J. Chen, W.D. Ren, X.S.Zhen, Corros. Sci. 49 (2007) 2436– 2449. [43] M.H. Shao, Y. Fu, R.G. Hu, C.J. Lin, Mater. Sci. Eng. A 344 (2003) 323–327. [44] K. Mizuno, A. Nylund, I. Olefjord, Corros. Sci. 43 (2001) 381–396.
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Figure captions Fig. 1 Hardness of the Al-Cu-Mg-Sc-Zr alloys aged at 190 oC. Fig. 2 Optical microstructures of the alloy in solid–solution state: (a) Al-Cu-Mg alloy; (b) Al-Cu-Mg-0.1Sc-0.2Zr alloy; (c) Al-Cu-Mg-0.2Sc-0.4Zr alloy. Fig.3 SEM-EBSD measurements of alloys peak-aged at 190oC. Orientation maps: (a) Al-Cu-Mg; (b) Al-Cu-Mg-0.1Sc-0.2Zr; (c) Al-Cu-Mg-0.2Sc-0.4Zr; Misorientation angle distribution: (d) Al-Cu-Mg, (e) Al-Cu-Mg-0.1Sc-0.2Zr, (f) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.4 TEM images of the studied alloys with different amount of Sc and Zr additions: (a) Al-Cu-Mg alloy; (b) Al-Cu-Mg-0.1Sc-0.2Zr alloy; (c) Al-Cu-Mg-0.2Sc-0.4Zr alloy; (d and e) Al3(Sc, Zr) in Al-Cu-Mg-0.2Sc-0.4Zr alloy; (f) SAD in [001]Al projection; (g) SAD in [011]Al projection. Fig.5 Metallographic images of cross-section morphology after immersion in the corrosive environment: (a) Al-Cu-Mg; (b) Al-Cu-Mg-0.1Sc-0.2Zr; (c) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.6 Metallographic images of cross-section morphology of the peak-aged alloys show the developed corrosion pits morphology after exposure to IGC environment for different durations: Al-Cu-Mg alloy exposed for (a) 3 h; (c) 6 h; (e) 12 h and (g) 24 h; Al-Cu-Mg-0.2Sc-0.4Zr alloy exposed for (b) 3 h; (d) 6 h; (f) 12 h and (h) 24 h; Schematic drawing of pit shape of (i)Al-Cu-Mg alloy; (j) Al-Cu-Mg-0.2Sc-0.4Zr. Fig.7 SEM images of (a) overall topography of the cross-section morphology; (b) high magnification of the cross-section morphology.
16
Figure List
Fig. 1
Fig. 2
Fig. 3
17
Fig. 4
18
Fig. 5
Fig. 6
Fig. 7
19
Table 1 Chemical compositions of studied alloys (in wt%), as determined by inductively coupled plasma–optical emission spectrometry (ICP–OES) Alloys
Cu
Mg
Zr
Sc
Fe
Si
Sc+Zr
Al
Al-Cu-Mg
3.92
2.11
0
0
0.17
0.06
0
Bal.
Al-Cu-Mg-0.1Sc-0.2Zr
4.12
1.89
0.23
0.09
0.16
0.07
0.32
Bal.
Al-Cu-Mg-0.2Sc-0.4Zr
3.96
1.86
0.41
0.18
0.17
0.07
0.59
Bal.
Table 2 Average EDS test results (three tests) of the grain boundary precipitates in Al-Cu-Mg-Sc-Zr alloys (wt%)
Element Al-Cu-Mg Al-Cu-Mg-0.1Sc-0.2Zr Al-Cu-Mg-0.2Sc-0.4Zr Cu
21.79
11.88
10.54
Mg
8.72
7.45
6.97
Cu:Mg
2.50
1.59
1.51
Table 3 Variations of the σb and δb with different amount of Sc and Zr addition for the peak-aged Al-Cu-Mg alloys
Alloys
Al-Cu-Mg
Al-Cu-Mg-0.1Sc-0.2Zr
Al-Cu-Mg-0.2Sc-0.4Zr
σb
σf
δUTS
δb
δf
δEI
(MPa)
(MPa)
(%)
(%)
(%)
(%)
363.5
320.3
11.88
7.36
2.64
64.13
(±17)
(±21)
(±0.1)
(±0.2)
436.0
411.5
13.64
10.32
(±14)
(±19)
(±0.2)
(±0.1)
432.5
412.6
14.16
12.36
(±18)
(±10)
(±0.1)
(±0.3)
20
5.64
4.60
24.34
12.71
Table 4 Average EDS test results (three tests) at the bottom of inter-granular corrosion of Al-Cu-Mg-Sc-Zr alloys after immersion for different time (wt%) Element
Al-Cu-Mg
Al-Cu-Mg-0.2Sc-0.4Zr
3h
3h
24 h
24 h
Cu
29.13 51.30
10.65
32.49
Mg
2.86
5.14
4.95
Cu:Mg
10.16 30.54
2.07
6.56
1.68
21