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Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys
Accepted Manuscript Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys Yaoyao Weng, Zhihong Jia, Lipeng Ding, Yanfeng Pan, Yingying Liu, Qing Liu PII:
S0925-8388(16)33604-0
DOI:
10.1016/j.jallcom.2016.11.140
Reference:
JALCOM 39636
To appear in:
Journal of Alloys and Compounds
Received Date: 22 September 2016 Revised Date:
31 October 2016
Accepted Date: 9 November 2016
Please cite this article as: Y. Weng, Z. Jia, L. Ding, Y. Pan, Y. Liu, Q. Liu, Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.140. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys Yaoyao Weng a, Zhihong Jia a, Lipeng Ding a, Yanfeng Pan b, Yingying Liu a, Qing Liu a College of Materials Science and Engineering, Chongqing University, Chongqing 400044 China
b
Suzhou Research Institute for Nonferrous Metals, No.200, Dongshenhu Road, Suzhou, 215026,
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a
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China
Abstract: The effects of Ag and/or Cu additions on the natural aging and precipitation hardening behavior of
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Al-Mg-Si alloys were investigated by using hardness test, differential scanning calorimetry (DSC), and transmission electron microscopy (TEM). Both Ag and Cu additions enhanced the hardening kinetics of Al-Mg-Si alloys during natural aging (NA) and artificial aging (AA) treatments. The strong interaction of Cu, Ag and Mg atoms is responsible for the improved precipitation kinetics of these alloys, resulting in a refinement of clusters and precipitates. The stronger interaction between Ag and Mg results in lower T4 hardness and higher AA
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hardness, compared to Cu. This is potentially beneficial for automobile applications where rapid hardening during paint baking is required. Both Ag and Cu additions reduce the detrimental effect of NA on subsequent AA, due to the formation of Ag or Cu-containing clusters. The clear advantage of Ag addition, compared to Cu, is the
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improvement of precipitation hardening response of Al-Mg-Si alloys.
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Keywords: Al-Mg-Si alloys; Ag addition; Cu addition; microstructure; precipitation hardening;
1. Introduction
Heat-treatable Al-Mg-Si (6xxx) alloys are becoming widely used for automobiles body panels due to a superior combination of high strength/weight ratio, good formability and corrosion resistance [1-4]. These alloys are strengthened by the precipitation of high density of nano-sized, semi-coherent or coherent metastable precipitates [5-7]. Al-Mg-Si alloy sheet for automobile production requires low yield strength and good formability in the T4 temper and high yield strength after a paint bake (PB) treatment for in-service dent resistance. Additions of other alloy elements to the ternary composition have been widely used to
ACCEPTED MANUSCRIPT improve the mechanical properties of 6xxx sheet. Cu additions to Al-Mg-Si alloys have been studied [8], effectively refining the precipitate distribution and enhancing precipitation kinetics during aging [9-11]. Cu suppresses the formation of β″ by the formation of other Cu-containing precipitates, such as L, S, C, Q', QP and QC [12-15]. The Q' phase and its precursors coexist with β″ contribute to an improved hardening
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capability. Another important effect of Cu additions, while not completely understood, is the reduction of the deleterious effect of prior NA on bake hardening. However, the relatively high yield strength of the high-Cu alloys in T4 temper can lead to inferior formability [16] and susceptibility to intergranular corrosion in Al-Mg-Si alloys.
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Small Ag additions to Al-Mg-Si alloys have been found to significantly stimulate precipitate formation, accelerating the aging kinetics and increasing peak hardness [17-19]. It has been assumed that Ag has a
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similar effect to Cu which creates a microstructure with a dense, uniform distribution of precipitates. Recent work has shown that Ag additions change the crystal structure of precipitates in Al-Mg-Si alloys. J. Nakamura and C.D. Marioara showed [18, 20, 21] that Ag enters the β' precipitate and replaces 1/3 of its Si atomic columns, changing its hexagonal unit cell parameter from 0.715nm to 0.690nm and the symmetry
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from space group P-63/m (β') to P-62m (β'Ag). The composition of this new β'Ag phase was identified as Al3Mg3Si2Ag, which coexists with the Ag-free β'-Mg9Si5 phase in the matrix. Silver also alters the structure of U2 phase by partly replacing Al, creating an expansion of the unit cell dimension from αU2=0.675nm to
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αU2Ag=0.690nm. Nakamura et al. [21] reported a strong silver enrichment at the interface of β'Ag/Al, at which most Al atom sites were replaced by Ag atoms. However, whether the Ag addition does fundamentally
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change the precipitation sequence of the Al-Mg-Si alloys is still unknown. K. Matsuda et al. [22] discovered a new quaternary grain boundary precipitate (Q'Ag) in an aged Al-Mg-Si-Ag alloy. The Q'Ag phase has a hexagonal crystal structure and a larger lattice parameter than that of Q'-phase in Al-Mg-Si-Cu alloys. However, the crystal structure of the Q'Ag phase has not been fully defined. S. Wenner et al. [23] investigated the distribution of Cu and Ag in Al-Mg-Si-Cu-Ag alloy precipitates. They found that Cu atoms located preferentially along coherent interfaces with the Al lattice. Ag localized at the narrow ends of precipitates, disrupting the local order of the precipitates. However, the combined effect of the Ag and Cu additions on mechanical properties of Al-Mg-Si alloys is still not completely understood.
ACCEPTED MANUSCRIPT While Ag and Cu additions to Al-Mg-Si alloys have been reported to create local symmetries and significantly improve the mechanical properties, the strengthening mechanism is still poorly understood. The individual influence of Ag or Cu on precipitation hardening in Al-Mg-Si alloys have been investigated in a number of studies, but relatively little work has addressed the effect of Ag and Cu in combination. Natural
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aging and its negative effect on AA strength in Ag/Cu-containing alloys are still not well understood. The present study has employed hardness testing, DSC, and TEM to study the influence of Ag and Cu on the microstructure and precipitation behavior of the Al-Mg-Si alloy.
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2. Experimental
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The compositions of the four experimental alloys are listed in Table 1. All these alloys have the same Mg and Si contents but different Ag and Cu additions (A2: 0.5 wt. % Ag-added, A3: 0.26 wt. % Ag-0.24 wt. % Cu-added, A4: 0.5 wt. % Cu-added). The four alloys were chill cast using industrially pure Al ingot and 99.9% purity Mg, Al-40 wt. % Si and Al-50 wt. % Cu, 99.9% purity Ag master alloys. The four alloys were homogenized at 560°C for 6 hours, followed by hot rolling and subsequent cold rolling to 1mm thick sheet.
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The cold rolled alloys were solution heat treated at 570°C for 20min and water quenched to room temperature (RT). Three types of aging treatment were applied, (1) NA at room temperature for different periods, (2) AA at 170°C for different times immediately after water quenching, (3) AA at 170°C for
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different times after two weeks of RT storage.
Vickers microhardness was measured after NA, AA and combined NA+AA treatments using a MH-5L
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microhardness tester with a load of 500g and a dwell time of 10s. The average of ten indentations distributed over the whole surface of each sample provided the data used to draw the hardness curves. The error in Vickers hardness is not more than ±3%. DSC was conducted under an argon atmosphere using a METTLER-1100LF system between 25ºC and 550ºC with a heating rate of 10ºC/min. TEM examination of the microstructures was done with a FEI Tecnai G2 F20 transmission electron microscope. Disc-shaped specimens for TEM were prepared by electro polishing using a Struers TenuPol-5 machine with an electrolyte of 1/3 HNO3 in methanol at a temperature of about -30°C. For each alloy condition, at least ten TEM bright-field images were taken to compare and make statistical comparisons.
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3. Results 3.1 Hardness measurements The hardness curves of the four alloys during RT storage are shown in Fig. 1 using a logarithmic time axis.
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The four alloys showed a rapid increase in hardness during the first two days, followed by a gentle increase up to 14 days. However, there are some differences between the reference alloy and the Ag- and/or Cu-containing alloys. Within the first two days NA, the Ag-containing alloy (A2) displayed similar hardness as the reference
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alloy (A1), suggesting that the effect of the individual addition of Ag on the hardness was not distinguishable at this stage. While the Cu addition could evidently decrease the hardness of the alloys (A3 and A4), and also the
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higher added Cu content the more hardness was decreased. After two days NA, the Ag- and/or Cu-containing alloys showed a gradually increase in hardness referring to the Al alloy, and up to 14 days NA, there was big gap of hardness between the Al alloy and the other three alloys. The hardness of the Ag-Cu-containing A3 alloy located in between the individual Ag-containing alloy and the individual Cu-containing alloy in the late stage of NA. This indicates that combined Ag and Cu additions lead to different clustering behavior.
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Fig. 2 is the hardness evolution of the four alloys aged at 170ºC for different aging time immediately after solution treatment and water quenching. All curves displayed similar profiles: the hardness increased rapidly to a maximum, followed by slightly decreasing and subsequently keeping as constant during over-aging stage. The
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Ag- and/or Cu-containing alloys exhibit enhanced precipitation hardening and overall higher AA hardness compared to the A1 base alloy. The times to reach peak hardness for the four alloys were 5h-A1, 2h-A2, 3h-A3
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and 4h-A4. The peak hardness of the Ag-containing A2 alloy was 133.3HV, approximately 7.6HV higher than the A1 base alloy (125.7HV), while the A3 and A4 alloys had similar and smaller peak hardness (131.2HV and 131.1HV, respectively). It can be seen that the addition of Ag is much more efficiently comparing to Cu in improving the precipitation kinetics in Al-Mg-Si alloys, particularly, in the early stage of AA, for instance, at 5min in this work. During over-aging stage, the hardness difference between the Ag and/or Cu added alloys became very smaller. The hardness curves of the four alloys artificially aged at 170ºC after two weeks RT storage are shown in Fig. 3. As expected, the hardness curves of the NA+AA samples showed overall lower hardness compared with the samples aged directly after quenching. The hardness values of the Ag- and/or Cu-containing alloys (A2, A3 and
ACCEPTED MANUSCRIPT A4) were all higher than that of the A1 base alloy. A small decrease in hardness was seen in the early stage of aging (about 0-5min), which is associated with the dissolution of clusters formed during NA. The times to reach peak hardness were longer (A1-6h, A2-4h. A3-5h and A4-6h) compared to that of the samples artificial aging immediately after quenching (the negative effect of prior NA). The precipitation hardening rate and the peak
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hardness of the Cu/Ag containing alloys were greatly enhanced compared to the A1 base alloy. The A3 (Ag+Cu) alloy had the highest peak hardness (119.1HV), while the A2 (Ag only) and A4 (Cu only) alloys displayed somewhat lower peak hardness (112.3HV for A2, 116.8HV for A4). The loss in the peak aged hardness (∆HV) of the four alloys due to the adverse effect of prior NA is given in Fig. 4. The A1 alloy has the largest ∆HV (27.8HV),
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showing the most serious negative effect of NA. The addition of Ag or Cu suppressed the detrimental effect of NA to various degrees. The A3 (Ag+Cu) alloy had the lowest ∆HV (12.1HV) and the A4 (Cu only) alloy had a
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slightly larger ∆HV (14.2HV), while the ∆HV of the A2 (Ag only) alloy was much larger (20.9HV). This indicates that Ag+Cu more effectively suppress the detrimental effect of NA compared to Ag or Cu only.
3.2 DSC analysis
DSC analysis was used to study the thermal events associated with precipitation of the four alloys. The DSC
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curves for the as-quenched alloys are given in Fig. 5(a). For the A1 base alloy, the exothermic peaks occurred can be correlated with the various stages of the precipitation process: exothermic peak I (81.7ºC, GP zones), exothermic peak II (266.5ºC, needle-like β″ precipitates), exothermic peak III (302.1ºC, β' precipitates), and
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exothermic peak IV (423.2ºC, formation of β and Si) [24]. The addition of Ag or Cu shifted the β″ exothermal peak to lower temperatures, implying an accelerated precipitation rate. The A2 (Ag only) alloy had a slightly
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higher β″ precipitation temperature and smaller peak area compared to that of the A4 (Cu only) alloy, which seems to contradict the higher precipitation kinetics of A2 alloy shown in Fig. 2. It is speculated that this is due to the small precipitation peak (exothermic peak V) located at 203.7ºC-231.3ºC in which the precipitation capability needed for β″ precipitation has been consumed partly. This small precipitation peak may relate to an Ag-containing pre-β″ phase formed in the early stage of aging. The DSC curves obtained after two weeks RT storage presented in Fig. 5(b) showed that the β″ precipitation peak was smaller and shifts to higher temperature compared to the as-quenched alloys, corresponding to the longer time needed to reach peak hardness. The β″ in the A3 (Ag+Cu) alloy had the smallest retarding effect (3.1oC), indicating the least negative effect of NA. The A2 and A4 alloys had similar small retarding effects (3.8oC and 4.1oC, respectively), while the A1 alloy showed the
ACCEPTED MANUSCRIPT largest adverse effect (6.4oC). These results indicate that the combined addition of Ag and Cu was most effective in suppressing the detrimental effect of prior NA, which is in agreement with the of the smallest hardness drop of the A3 alloy shown in Fig. 4.
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3.3. Microstructure investigation Fig. 6 shows TEM bright-field images of the four alloys peak aged at 170°C immediately after quenching. All images were acquired in <100>Al directions in this work. Large numbers of needle-like precipitates were observed in all four alloys. A high density of lath-like precipitates was also seen in the A3 (Ag+Cu) and A4 (Cu
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only) alloys. From the HRTEM images and the corresponding Fast Fourier Filtering transforms (FFT) patterns shown in Fig.7, the needle-like precipitates were identified as the β″ phase and the lath-like precipitates with the
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cross-section elongated along <510>Al were identified as the Q' phase. The lath-like precipitates with the cross-section elongated along <100>Al were the L phase, a precursor of Q' with a significant strengthening effect. The coexistence of lath-like precipitates (Q' and L) with needle-like β″ implies that the addition of Cu has suppressed the formation of β″ by the formation of Cu-containing precipitates. The precipitate length distributions (PLD) of the four alloys peak aged at 170 oC immediately after water quenching are plotted in Fig. 9. The A2 (Ag
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only) alloy had narrower precipitates distribution with an average length of 6.099nm compared to the other alloys, implying that the effect of Ag on refining the precipitates is more effective than Cu. Fig.8 shows TEM bright-field images of the four alloys peak-aged at 170 oC after two weeks RT storage. The
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type of precipitates was not changed compare to that of alloys peak-aged immediately after quenching, while the precipitates became larger and the number density decreased in each alloy condition. The precipitate length
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distributions (PLD) of the four alloys peak aged at 170 oC after two weeks of natural aging is plotted in Fig. 10. The distribution of precipitates was very narrow in the direct peak aged AA condition of A2 alloy, however, after NA for 2 weeks before AA, the length distribution was broader and the average length increased to 29.5nm. The PLD for the A3 (Ag+Cu) and A4 (Cu only) alloys also experienced a similar evolution. The average lengths for these two alloys increased from 5.9nm and 8.3nm to 21.1nm and 26.5nm, respectively. It can be seen that the A2 (Ag only) alloy showed a distinct length increase (23.5nm) compared to the Cu-containing alloys (15.2nm for A3 and 18.3nm for A4 alloy), implying that the Cu addition is more effectively in suppressing the negative effect of NA compared to the Ag addition.
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4. Discussion From the hardness curves shown in Fig. 1, it is evident that Cu addition results in lower hardness during the
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early stages of NA and higher hardness after a long time at RT compared to the A1 base alloy. While the A2 (Ag only) alloy displayed similar hardness as the reference alloy (A1) within two days of NA. It is known that natural
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aging is a complex interaction between vacancies and solute atoms. The presence of quenched-in vacancies plays an important role in controlling clustering kinetics during NA [25-27]. For Al-Mg-Si alloys, a high density of
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quenched-in vacancies is present and they are quickly trapped by solute atoms to form clusters after the quench. The Si-rich clusters are dominant in this stage due to the high diffusivity and low solubility of Si atoms in the Al matrix. Then Mg atoms slowly diffuse into these clusters over at least one week and eventually form Mg-Si co-clusters. During natural aging, solute atoms and vacancies are consumed to form clusters, and the rate of
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clustering decreases as NA proceeds. For the Cu/Ag-containing alloys, Cu atoms can increase the free energy of quenching clusters in the initial clustering stage of NA [28]. Incorporation of Cu atoms into the clusters
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minimizes the strain energy caused by the difference in atomic sizes, and makes the clusters more stable [29]. Thus, the formation of quenching clusters is suppressed in Cu-containing alloys, leading to a lower hardness in
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early stages of NA. However, for longer RT storage times, the large supersaturation in Cu-containing alloys produces a larger driving force for clustering. Therefore, the additions of Cu tend to accelerate the kinetics in later stages of NA. The Ag addition has a similar effect to that of Cu, increasing the free energy of quenching clusters to retard the formation of clusters in the early stage of NA. The Ag atom has a larger radius than the Cu atom, making it less effective in minimizing the strain energy of clusters. This leads to a lower driving force compared
with Cu atom after a long period of RT storage. Therefore, the Ag-containing A2 alloy has a lower T4 temper hardness than the Cu-containing A4 alloy. A relative lower T4 hardness of the A3 (Ag+Cu) alloy
ACCEPTED MANUSCRIPT than the A4 (Cu only) alloy due to the low Cu content. Both Ag and Cu additions increase precipitation kinetics and refine the precipitate distribution of Al-Mg-Si alloys during artificial aging immediately after quenching. G. A. Edwards et al. [24] investigated
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the evolution of clusters using atom probe field ion microscope (APFIM). They proposed that the initial clustering sequence is as follow: Al-SSS → clusters of Si atoms and clusters of Mg atoms → dissolution of Mg clusters → formation of Mg-Si co-clusters. The addition of Cu and/or Ag, significantly changes the
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clustering during the initial stages of AA. As shown in Fig.11 [30], Cu atoms have a strong interaction with
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Mg atoms, which can partly take away the Mg atoms from Mg-Si co-clusters to form the Mg-Si-Cu co-clusters, leading to a finer distribution of clusters. For Cu-containing alloys, Cu-rich clusters evolve into Q' and L, and the Cu-free clusters evolve into β″ phase. The coexistence of Q', L and β″ phases contribute to the high peak aged hardness of these alloys. Since Ag atoms have a stronger binding energy with Mg atoms
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and vacancies compared to Cu atoms, the addition of Ag is more effectively in refining the cluster distribution, which results in a higher precipitation kinetics and precipitation hardening response. No other precipitates were found in the Ag-containing A2 alloy, revealing that the refined β″ distribution was
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responsible for the high peak aged hardness value.
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It has long been known that prior natural aging can significantly suppress the precipitation kinetics of artificial aging in Al-Mg-Si alloys. Ag and Cu additions can reduce the negative effect of NA, especially combination. It could be deduced that the different detrimental effect of NA may associated with the clusters formed during NA. A. Serizawa et al. [31] reported the clusters formed during NA had a wide range of Mg/Si ratio and did not readily transform to β″. The formation of these NA clusters retards the precipitation of β″ during AA because the quenched-in vacancies and solute elements were consumed. Besides, these NA clusters are thermally stable and hard to be dissolved at the artificial aging temperature. The addition of other alloying elements can lead to significant changes in the composition of these NA clusters. Cu (or Ag) atoms can incorporated into the Mg-Si co-clusters and change the stability of these clusters during NA, making
ACCEPTED MANUSCRIPT them relatively easier to transform to β″ phase during artificial aging. However, the underlying mechanism for why the Cu is more effective in suppressing the negative effect of NA compared to Ag is still unclear. It is supposed that the stability of the Ag-containing clusters formed during NA is greater than that of the Cu-containing clusters, retarding the transformation of these Ag-containing clusters to β″ during AA. The
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combined addition of Ag and Cu is more effective in suppressing the adverse effect of prior NA compared to the individual addition of Ag or Cu, which may be of interest for industrial production.
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5. Conclusions
The effects of Ag and Cu additions on the natural aging, artificial aging and the negative effect of NA
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were systematically investigated. The results are summarized as follows.
(1) The addition of Ag does not influence hardness within first two days of NA, but improves hardness in the late stage of NA. While the addition of Cu results in lower hardness during the early stages of NA (within two days) and higher hardness after a longer period of RT storage (two weeks). The difference between Ag and Cu additions could be explained by distinguishable interaction of Ag or Cu atoms with
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clusters.
(2) During artificial aging, the addition of Cu or Ag increases the precipitation kinetics and refines the precipitate distribution of Al-Mg-Si alloys. The A2 (Ag only) alloy has a higher age-hardening response than
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that of A4 (Cu only) alloy due to the stronger binding energy between Ag and Mg atoms and vacancies compared to that of Cu atoms.
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(3) The negative effect of NA on subsequent AA is decreased by the addition of Ag or Cu, especially combination. Cu additions are more effective in suppressing the adverse effect of natural aging compared with Ag.
Acknowledgements The authors wish to acknowledge Prof. Robert E. Sanders from Chongqing University for English writing improvements and comments. This work was supported by International Science & Technology Cooperation Program of China (Grant No. 2014DFA51270), National Natural Science Foundation of China
ACCEPTED MANUSCRIPT (Grant No. 51271209) and the Foundation for Innovative Research Groups of the National Natural Science Foundation of China (Grant No. 51421001).
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Figure caption Fig. 1 Evolution of Vickers hardness of four alloys during NA after water quenching. Fig. 2. Evolution of Vickers hardness of four alloys during AA at 170°C immediately after
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quenching. Fig. 3. Evolution of Vickers hardness of four alloys during AA at 170°C after two weeks of RT
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storage.
Fig. 4. The hardness dropped (∆HV) induced by NA of the four alloys in the peak aging condition.
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Fig. 5. DSC heat flow curves of four alloys. (a) immediately after quenching, (b) after two weeks of NA.
Fig. 6. TEM bright-field images of artificial aging at 170°C immediately after quenching for four alloys: (a) A1 alloy, the base alloy; (b) A2 alloy, with 0.5wt. % Ag; (c) A3 alloy, with 0.25wt. %
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Ag and 0.25wt. % Cu; (d) A4 alloy, with 0.5wt. % Cu. Fig. 7. HRTEM images and corresponding FFT patterns of different type of precipitates. (a and b) β″
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precipitate, (c and d) Q’ precipitate, (e and f) L phase.
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Fig. 8. TEM bright-field images of artificial aging at 170°C after two weeks of NA for four alloys: (a) A1 alloy, the base alloy; (b) A2 alloy, with 0.5wt. % Ag; (c) A3 alloy, with 0.25wt. % Ag and 0.25wt. % Cu; (d) A4 alloy, with 0.5wt. % Cu. Fig. 9. Average length and length distribution of needle-like βʺ precipitates in the AA peak aged samples shown in Fig. 6. Fig. 10. Average length and length distribution of needle-like βʺ precipitates in the NA+AA peak aged samples shown in Fig. 8.
ACCEPTED MANUSCRIPT Fig. 11. First-principles calculated nearest-neighbor solute–□ binding energies as a function of the
Table caption
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Table 1 Compositions of the four alloys (wt. %).
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solute volume and interaction energy map of Al-Mg-Si-X alloys (X: microalloying element) [30].
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Table 1 Compositions of the four alloys (wt. %). Fe 0.11 0.11 0.11 0.13
Mn 0.062 0.06 0.056 0.06
Ag 0.49 0.26 -
Cu 0.24 0.5
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Mg 1.09 1.13 1.14 1.1
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ACCEPTED MANUSCRIPT Highlights The Ag-containing alloy had relatively lower T4 hardness compared to the Cu-containing alloy.
Ag addition improves the precipitation kinetics during AA for the strong interaction between Ag, Mg and vacancy.
The negative effect of NA is decreased by the addition of Ag or Cu, especially combination.
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S0925-8388(16)33604-0
DOI:
10.1016/j.jallcom.2016.11.140
Reference:
JALCOM 39636
To appear in:
Journal of Alloys and Compounds
Received Date: 22 September 2016 Revised Date:
31 October 2016
Accepted Date: 9 November 2016
Please cite this article as: Y. Weng, Z. Jia, L. Ding, Y. Pan, Y. Liu, Q. Liu, Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.11.140. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Effect of Ag and Cu additions on natural aging and precipitation hardening behavior in Al-Mg-Si alloys Yaoyao Weng a, Zhihong Jia a, Lipeng Ding a, Yanfeng Pan b, Yingying Liu a, Qing Liu a College of Materials Science and Engineering, Chongqing University, Chongqing 400044 China
b
Suzhou Research Institute for Nonferrous Metals, No.200, Dongshenhu Road, Suzhou, 215026,
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a
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China
Abstract: The effects of Ag and/or Cu additions on the natural aging and precipitation hardening behavior of
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Al-Mg-Si alloys were investigated by using hardness test, differential scanning calorimetry (DSC), and transmission electron microscopy (TEM). Both Ag and Cu additions enhanced the hardening kinetics of Al-Mg-Si alloys during natural aging (NA) and artificial aging (AA) treatments. The strong interaction of Cu, Ag and Mg atoms is responsible for the improved precipitation kinetics of these alloys, resulting in a refinement of clusters and precipitates. The stronger interaction between Ag and Mg results in lower T4 hardness and higher AA
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hardness, compared to Cu. This is potentially beneficial for automobile applications where rapid hardening during paint baking is required. Both Ag and Cu additions reduce the detrimental effect of NA on subsequent AA, due to the formation of Ag or Cu-containing clusters. The clear advantage of Ag addition, compared to Cu, is the
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improvement of precipitation hardening response of Al-Mg-Si alloys.
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Keywords: Al-Mg-Si alloys; Ag addition; Cu addition; microstructure; precipitation hardening;
1. Introduction
Heat-treatable Al-Mg-Si (6xxx) alloys are becoming widely used for automobiles body panels due to a superior combination of high strength/weight ratio, good formability and corrosion resistance [1-4]. These alloys are strengthened by the precipitation of high density of nano-sized, semi-coherent or coherent metastable precipitates [5-7]. Al-Mg-Si alloy sheet for automobile production requires low yield strength and good formability in the T4 temper and high yield strength after a paint bake (PB) treatment for in-service dent resistance. Additions of other alloy elements to the ternary composition have been widely used to
ACCEPTED MANUSCRIPT improve the mechanical properties of 6xxx sheet. Cu additions to Al-Mg-Si alloys have been studied [8], effectively refining the precipitate distribution and enhancing precipitation kinetics during aging [9-11]. Cu suppresses the formation of β″ by the formation of other Cu-containing precipitates, such as L, S, C, Q', QP and QC [12-15]. The Q' phase and its precursors coexist with β″ contribute to an improved hardening
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capability. Another important effect of Cu additions, while not completely understood, is the reduction of the deleterious effect of prior NA on bake hardening. However, the relatively high yield strength of the high-Cu alloys in T4 temper can lead to inferior formability [16] and susceptibility to intergranular corrosion in Al-Mg-Si alloys.
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Small Ag additions to Al-Mg-Si alloys have been found to significantly stimulate precipitate formation, accelerating the aging kinetics and increasing peak hardness [17-19]. It has been assumed that Ag has a
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similar effect to Cu which creates a microstructure with a dense, uniform distribution of precipitates. Recent work has shown that Ag additions change the crystal structure of precipitates in Al-Mg-Si alloys. J. Nakamura and C.D. Marioara showed [18, 20, 21] that Ag enters the β' precipitate and replaces 1/3 of its Si atomic columns, changing its hexagonal unit cell parameter from 0.715nm to 0.690nm and the symmetry
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from space group P-63/m (β') to P-62m (β'Ag). The composition of this new β'Ag phase was identified as Al3Mg3Si2Ag, which coexists with the Ag-free β'-Mg9Si5 phase in the matrix. Silver also alters the structure of U2 phase by partly replacing Al, creating an expansion of the unit cell dimension from αU2=0.675nm to
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αU2Ag=0.690nm. Nakamura et al. [21] reported a strong silver enrichment at the interface of β'Ag/Al, at which most Al atom sites were replaced by Ag atoms. However, whether the Ag addition does fundamentally
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change the precipitation sequence of the Al-Mg-Si alloys is still unknown. K. Matsuda et al. [22] discovered a new quaternary grain boundary precipitate (Q'Ag) in an aged Al-Mg-Si-Ag alloy. The Q'Ag phase has a hexagonal crystal structure and a larger lattice parameter than that of Q'-phase in Al-Mg-Si-Cu alloys. However, the crystal structure of the Q'Ag phase has not been fully defined. S. Wenner et al. [23] investigated the distribution of Cu and Ag in Al-Mg-Si-Cu-Ag alloy precipitates. They found that Cu atoms located preferentially along coherent interfaces with the Al lattice. Ag localized at the narrow ends of precipitates, disrupting the local order of the precipitates. However, the combined effect of the Ag and Cu additions on mechanical properties of Al-Mg-Si alloys is still not completely understood.
ACCEPTED MANUSCRIPT While Ag and Cu additions to Al-Mg-Si alloys have been reported to create local symmetries and significantly improve the mechanical properties, the strengthening mechanism is still poorly understood. The individual influence of Ag or Cu on precipitation hardening in Al-Mg-Si alloys have been investigated in a number of studies, but relatively little work has addressed the effect of Ag and Cu in combination. Natural
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aging and its negative effect on AA strength in Ag/Cu-containing alloys are still not well understood. The present study has employed hardness testing, DSC, and TEM to study the influence of Ag and Cu on the microstructure and precipitation behavior of the Al-Mg-Si alloy.
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2. Experimental
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The compositions of the four experimental alloys are listed in Table 1. All these alloys have the same Mg and Si contents but different Ag and Cu additions (A2: 0.5 wt. % Ag-added, A3: 0.26 wt. % Ag-0.24 wt. % Cu-added, A4: 0.5 wt. % Cu-added). The four alloys were chill cast using industrially pure Al ingot and 99.9% purity Mg, Al-40 wt. % Si and Al-50 wt. % Cu, 99.9% purity Ag master alloys. The four alloys were homogenized at 560°C for 6 hours, followed by hot rolling and subsequent cold rolling to 1mm thick sheet.
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The cold rolled alloys were solution heat treated at 570°C for 20min and water quenched to room temperature (RT). Three types of aging treatment were applied, (1) NA at room temperature for different periods, (2) AA at 170°C for different times immediately after water quenching, (3) AA at 170°C for
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different times after two weeks of RT storage.
Vickers microhardness was measured after NA, AA and combined NA+AA treatments using a MH-5L
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microhardness tester with a load of 500g and a dwell time of 10s. The average of ten indentations distributed over the whole surface of each sample provided the data used to draw the hardness curves. The error in Vickers hardness is not more than ±3%. DSC was conducted under an argon atmosphere using a METTLER-1100LF system between 25ºC and 550ºC with a heating rate of 10ºC/min. TEM examination of the microstructures was done with a FEI Tecnai G2 F20 transmission electron microscope. Disc-shaped specimens for TEM were prepared by electro polishing using a Struers TenuPol-5 machine with an electrolyte of 1/3 HNO3 in methanol at a temperature of about -30°C. For each alloy condition, at least ten TEM bright-field images were taken to compare and make statistical comparisons.
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3. Results 3.1 Hardness measurements The hardness curves of the four alloys during RT storage are shown in Fig. 1 using a logarithmic time axis.
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The four alloys showed a rapid increase in hardness during the first two days, followed by a gentle increase up to 14 days. However, there are some differences between the reference alloy and the Ag- and/or Cu-containing alloys. Within the first two days NA, the Ag-containing alloy (A2) displayed similar hardness as the reference
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alloy (A1), suggesting that the effect of the individual addition of Ag on the hardness was not distinguishable at this stage. While the Cu addition could evidently decrease the hardness of the alloys (A3 and A4), and also the
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higher added Cu content the more hardness was decreased. After two days NA, the Ag- and/or Cu-containing alloys showed a gradually increase in hardness referring to the Al alloy, and up to 14 days NA, there was big gap of hardness between the Al alloy and the other three alloys. The hardness of the Ag-Cu-containing A3 alloy located in between the individual Ag-containing alloy and the individual Cu-containing alloy in the late stage of NA. This indicates that combined Ag and Cu additions lead to different clustering behavior.
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Fig. 2 is the hardness evolution of the four alloys aged at 170ºC for different aging time immediately after solution treatment and water quenching. All curves displayed similar profiles: the hardness increased rapidly to a maximum, followed by slightly decreasing and subsequently keeping as constant during over-aging stage. The
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Ag- and/or Cu-containing alloys exhibit enhanced precipitation hardening and overall higher AA hardness compared to the A1 base alloy. The times to reach peak hardness for the four alloys were 5h-A1, 2h-A2, 3h-A3
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and 4h-A4. The peak hardness of the Ag-containing A2 alloy was 133.3HV, approximately 7.6HV higher than the A1 base alloy (125.7HV), while the A3 and A4 alloys had similar and smaller peak hardness (131.2HV and 131.1HV, respectively). It can be seen that the addition of Ag is much more efficiently comparing to Cu in improving the precipitation kinetics in Al-Mg-Si alloys, particularly, in the early stage of AA, for instance, at 5min in this work. During over-aging stage, the hardness difference between the Ag and/or Cu added alloys became very smaller. The hardness curves of the four alloys artificially aged at 170ºC after two weeks RT storage are shown in Fig. 3. As expected, the hardness curves of the NA+AA samples showed overall lower hardness compared with the samples aged directly after quenching. The hardness values of the Ag- and/or Cu-containing alloys (A2, A3 and
ACCEPTED MANUSCRIPT A4) were all higher than that of the A1 base alloy. A small decrease in hardness was seen in the early stage of aging (about 0-5min), which is associated with the dissolution of clusters formed during NA. The times to reach peak hardness were longer (A1-6h, A2-4h. A3-5h and A4-6h) compared to that of the samples artificial aging immediately after quenching (the negative effect of prior NA). The precipitation hardening rate and the peak
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hardness of the Cu/Ag containing alloys were greatly enhanced compared to the A1 base alloy. The A3 (Ag+Cu) alloy had the highest peak hardness (119.1HV), while the A2 (Ag only) and A4 (Cu only) alloys displayed somewhat lower peak hardness (112.3HV for A2, 116.8HV for A4). The loss in the peak aged hardness (∆HV) of the four alloys due to the adverse effect of prior NA is given in Fig. 4. The A1 alloy has the largest ∆HV (27.8HV),
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showing the most serious negative effect of NA. The addition of Ag or Cu suppressed the detrimental effect of NA to various degrees. The A3 (Ag+Cu) alloy had the lowest ∆HV (12.1HV) and the A4 (Cu only) alloy had a
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slightly larger ∆HV (14.2HV), while the ∆HV of the A2 (Ag only) alloy was much larger (20.9HV). This indicates that Ag+Cu more effectively suppress the detrimental effect of NA compared to Ag or Cu only.
3.2 DSC analysis
DSC analysis was used to study the thermal events associated with precipitation of the four alloys. The DSC
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curves for the as-quenched alloys are given in Fig. 5(a). For the A1 base alloy, the exothermic peaks occurred can be correlated with the various stages of the precipitation process: exothermic peak I (81.7ºC, GP zones), exothermic peak II (266.5ºC, needle-like β″ precipitates), exothermic peak III (302.1ºC, β' precipitates), and
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exothermic peak IV (423.2ºC, formation of β and Si) [24]. The addition of Ag or Cu shifted the β″ exothermal peak to lower temperatures, implying an accelerated precipitation rate. The A2 (Ag only) alloy had a slightly
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higher β″ precipitation temperature and smaller peak area compared to that of the A4 (Cu only) alloy, which seems to contradict the higher precipitation kinetics of A2 alloy shown in Fig. 2. It is speculated that this is due to the small precipitation peak (exothermic peak V) located at 203.7ºC-231.3ºC in which the precipitation capability needed for β″ precipitation has been consumed partly. This small precipitation peak may relate to an Ag-containing pre-β″ phase formed in the early stage of aging. The DSC curves obtained after two weeks RT storage presented in Fig. 5(b) showed that the β″ precipitation peak was smaller and shifts to higher temperature compared to the as-quenched alloys, corresponding to the longer time needed to reach peak hardness. The β″ in the A3 (Ag+Cu) alloy had the smallest retarding effect (3.1oC), indicating the least negative effect of NA. The A2 and A4 alloys had similar small retarding effects (3.8oC and 4.1oC, respectively), while the A1 alloy showed the
ACCEPTED MANUSCRIPT largest adverse effect (6.4oC). These results indicate that the combined addition of Ag and Cu was most effective in suppressing the detrimental effect of prior NA, which is in agreement with the of the smallest hardness drop of the A3 alloy shown in Fig. 4.
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3.3. Microstructure investigation Fig. 6 shows TEM bright-field images of the four alloys peak aged at 170°C immediately after quenching. All images were acquired in <100>Al directions in this work. Large numbers of needle-like precipitates were observed in all four alloys. A high density of lath-like precipitates was also seen in the A3 (Ag+Cu) and A4 (Cu
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only) alloys. From the HRTEM images and the corresponding Fast Fourier Filtering transforms (FFT) patterns shown in Fig.7, the needle-like precipitates were identified as the β″ phase and the lath-like precipitates with the
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cross-section elongated along <510>Al were identified as the Q' phase. The lath-like precipitates with the cross-section elongated along <100>Al were the L phase, a precursor of Q' with a significant strengthening effect. The coexistence of lath-like precipitates (Q' and L) with needle-like β″ implies that the addition of Cu has suppressed the formation of β″ by the formation of Cu-containing precipitates. The precipitate length distributions (PLD) of the four alloys peak aged at 170 oC immediately after water quenching are plotted in Fig. 9. The A2 (Ag
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only) alloy had narrower precipitates distribution with an average length of 6.099nm compared to the other alloys, implying that the effect of Ag on refining the precipitates is more effective than Cu. Fig.8 shows TEM bright-field images of the four alloys peak-aged at 170 oC after two weeks RT storage. The
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type of precipitates was not changed compare to that of alloys peak-aged immediately after quenching, while the precipitates became larger and the number density decreased in each alloy condition. The precipitate length
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distributions (PLD) of the four alloys peak aged at 170 oC after two weeks of natural aging is plotted in Fig. 10. The distribution of precipitates was very narrow in the direct peak aged AA condition of A2 alloy, however, after NA for 2 weeks before AA, the length distribution was broader and the average length increased to 29.5nm. The PLD for the A3 (Ag+Cu) and A4 (Cu only) alloys also experienced a similar evolution. The average lengths for these two alloys increased from 5.9nm and 8.3nm to 21.1nm and 26.5nm, respectively. It can be seen that the A2 (Ag only) alloy showed a distinct length increase (23.5nm) compared to the Cu-containing alloys (15.2nm for A3 and 18.3nm for A4 alloy), implying that the Cu addition is more effectively in suppressing the negative effect of NA compared to the Ag addition.
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4. Discussion From the hardness curves shown in Fig. 1, it is evident that Cu addition results in lower hardness during the
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early stages of NA and higher hardness after a long time at RT compared to the A1 base alloy. While the A2 (Ag only) alloy displayed similar hardness as the reference alloy (A1) within two days of NA. It is known that natural
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aging is a complex interaction between vacancies and solute atoms. The presence of quenched-in vacancies plays an important role in controlling clustering kinetics during NA [25-27]. For Al-Mg-Si alloys, a high density of
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quenched-in vacancies is present and they are quickly trapped by solute atoms to form clusters after the quench. The Si-rich clusters are dominant in this stage due to the high diffusivity and low solubility of Si atoms in the Al matrix. Then Mg atoms slowly diffuse into these clusters over at least one week and eventually form Mg-Si co-clusters. During natural aging, solute atoms and vacancies are consumed to form clusters, and the rate of
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clustering decreases as NA proceeds. For the Cu/Ag-containing alloys, Cu atoms can increase the free energy of quenching clusters in the initial clustering stage of NA [28]. Incorporation of Cu atoms into the clusters
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minimizes the strain energy caused by the difference in atomic sizes, and makes the clusters more stable [29]. Thus, the formation of quenching clusters is suppressed in Cu-containing alloys, leading to a lower hardness in
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early stages of NA. However, for longer RT storage times, the large supersaturation in Cu-containing alloys produces a larger driving force for clustering. Therefore, the additions of Cu tend to accelerate the kinetics in later stages of NA. The Ag addition has a similar effect to that of Cu, increasing the free energy of quenching clusters to retard the formation of clusters in the early stage of NA. The Ag atom has a larger radius than the Cu atom, making it less effective in minimizing the strain energy of clusters. This leads to a lower driving force compared
with Cu atom after a long period of RT storage. Therefore, the Ag-containing A2 alloy has a lower T4 temper hardness than the Cu-containing A4 alloy. A relative lower T4 hardness of the A3 (Ag+Cu) alloy
ACCEPTED MANUSCRIPT than the A4 (Cu only) alloy due to the low Cu content. Both Ag and Cu additions increase precipitation kinetics and refine the precipitate distribution of Al-Mg-Si alloys during artificial aging immediately after quenching. G. A. Edwards et al. [24] investigated
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the evolution of clusters using atom probe field ion microscope (APFIM). They proposed that the initial clustering sequence is as follow: Al-SSS → clusters of Si atoms and clusters of Mg atoms → dissolution of Mg clusters → formation of Mg-Si co-clusters. The addition of Cu and/or Ag, significantly changes the
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clustering during the initial stages of AA. As shown in Fig.11 [30], Cu atoms have a strong interaction with
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Mg atoms, which can partly take away the Mg atoms from Mg-Si co-clusters to form the Mg-Si-Cu co-clusters, leading to a finer distribution of clusters. For Cu-containing alloys, Cu-rich clusters evolve into Q' and L, and the Cu-free clusters evolve into β″ phase. The coexistence of Q', L and β″ phases contribute to the high peak aged hardness of these alloys. Since Ag atoms have a stronger binding energy with Mg atoms
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and vacancies compared to Cu atoms, the addition of Ag is more effectively in refining the cluster distribution, which results in a higher precipitation kinetics and precipitation hardening response. No other precipitates were found in the Ag-containing A2 alloy, revealing that the refined β″ distribution was
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responsible for the high peak aged hardness value.
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It has long been known that prior natural aging can significantly suppress the precipitation kinetics of artificial aging in Al-Mg-Si alloys. Ag and Cu additions can reduce the negative effect of NA, especially combination. It could be deduced that the different detrimental effect of NA may associated with the clusters formed during NA. A. Serizawa et al. [31] reported the clusters formed during NA had a wide range of Mg/Si ratio and did not readily transform to β″. The formation of these NA clusters retards the precipitation of β″ during AA because the quenched-in vacancies and solute elements were consumed. Besides, these NA clusters are thermally stable and hard to be dissolved at the artificial aging temperature. The addition of other alloying elements can lead to significant changes in the composition of these NA clusters. Cu (or Ag) atoms can incorporated into the Mg-Si co-clusters and change the stability of these clusters during NA, making
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combined addition of Ag and Cu is more effective in suppressing the adverse effect of prior NA compared to the individual addition of Ag or Cu, which may be of interest for industrial production.
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5. Conclusions
The effects of Ag and Cu additions on the natural aging, artificial aging and the negative effect of NA
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were systematically investigated. The results are summarized as follows.
(1) The addition of Ag does not influence hardness within first two days of NA, but improves hardness in the late stage of NA. While the addition of Cu results in lower hardness during the early stages of NA (within two days) and higher hardness after a longer period of RT storage (two weeks). The difference between Ag and Cu additions could be explained by distinguishable interaction of Ag or Cu atoms with
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clusters.
(2) During artificial aging, the addition of Cu or Ag increases the precipitation kinetics and refines the precipitate distribution of Al-Mg-Si alloys. The A2 (Ag only) alloy has a higher age-hardening response than
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that of A4 (Cu only) alloy due to the stronger binding energy between Ag and Mg atoms and vacancies compared to that of Cu atoms.
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(3) The negative effect of NA on subsequent AA is decreased by the addition of Ag or Cu, especially combination. Cu additions are more effective in suppressing the adverse effect of natural aging compared with Ag.
Acknowledgements The authors wish to acknowledge Prof. Robert E. Sanders from Chongqing University for English writing improvements and comments. This work was supported by International Science & Technology Cooperation Program of China (Grant No. 2014DFA51270), National Natural Science Foundation of China
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Figure caption Fig. 1 Evolution of Vickers hardness of four alloys during NA after water quenching. Fig. 2. Evolution of Vickers hardness of four alloys during AA at 170°C immediately after
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quenching. Fig. 3. Evolution of Vickers hardness of four alloys during AA at 170°C after two weeks of RT
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storage.
Fig. 4. The hardness dropped (∆HV) induced by NA of the four alloys in the peak aging condition.
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Fig. 5. DSC heat flow curves of four alloys. (a) immediately after quenching, (b) after two weeks of NA.
Fig. 6. TEM bright-field images of artificial aging at 170°C immediately after quenching for four alloys: (a) A1 alloy, the base alloy; (b) A2 alloy, with 0.5wt. % Ag; (c) A3 alloy, with 0.25wt. %
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Ag and 0.25wt. % Cu; (d) A4 alloy, with 0.5wt. % Cu. Fig. 7. HRTEM images and corresponding FFT patterns of different type of precipitates. (a and b) β″
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precipitate, (c and d) Q’ precipitate, (e and f) L phase.
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Fig. 8. TEM bright-field images of artificial aging at 170°C after two weeks of NA for four alloys: (a) A1 alloy, the base alloy; (b) A2 alloy, with 0.5wt. % Ag; (c) A3 alloy, with 0.25wt. % Ag and 0.25wt. % Cu; (d) A4 alloy, with 0.5wt. % Cu. Fig. 9. Average length and length distribution of needle-like βʺ precipitates in the AA peak aged samples shown in Fig. 6. Fig. 10. Average length and length distribution of needle-like βʺ precipitates in the NA+AA peak aged samples shown in Fig. 8.
ACCEPTED MANUSCRIPT Fig. 11. First-principles calculated nearest-neighbor solute–□ binding energies as a function of the
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Table 1 Compositions of the four alloys (wt. %).
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solute volume and interaction energy map of Al-Mg-Si-X alloys (X: microalloying element) [30].
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Table 1 Compositions of the four alloys (wt. %). Fe 0.11 0.11 0.11 0.13
Mn 0.062 0.06 0.056 0.06
Ag 0.49 0.26 -
Cu 0.24 0.5
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Mg 1.09 1.13 1.14 1.1
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A1 A2 A3 A4
Si 0.71 0.71 0.71 0.70
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ACCEPTED MANUSCRIPT Highlights The Ag-containing alloy had relatively lower T4 hardness compared to the Cu-containing alloy.
Ag addition improves the precipitation kinetics during AA for the strong interaction between Ag, Mg and vacancy.
The negative effect of NA is decreased by the addition of Ag or Cu, especially combination.
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