Microstructure, mechanical properties and stress corrosion cracking of Al–Zn–Mg–Zr alloy sheet with trace amount of Sc

Accepted Manuscript Microstructure, mechanical properties and stress corrosion cracking of Al-Zn-Mg-Zr alloy sheet with trace amount of Sc Xing Huang, Qinglin Pan, Bo Li, Zhiming Liu, Zhiqi Huang, Zhimin Yin PII:

S0925-8388(15)30706-4

DOI:

10.1016/j.jallcom.2015.08.011

Reference:

JALCOM 35005

To appear in:

Journal of Alloys and Compounds

Received Date: 17 April 2015 Revised Date:

19 July 2015

Accepted Date: 1 August 2015

Please cite this article as: X. Huang, Q. Pan, B. Li, Z. Liu, Z. Huang, Z. Yin, Microstructure, mechanical properties and stress corrosion cracking of Al-Zn-Mg-Zr alloy sheet with trace amount of Sc, Journal of Alloys and Compounds (2015), doi: 10.1016/j.jallcom.2015.08.011. 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.

ACCEPTED MANUSCRIPT

Microstructure, mechanical properties and stress corrosion cracking of Al-Zn-Mg-Zr alloy sheet with trace amount of Sc

RI PT

Xing Huanga, Qinglin Pana, *, Bo Lia, Zhiming Liub, Zhiqi Huangb, Zhimin Yina a

School of Materials Science and Engineering, Central South University, Changsha 410083,

Guangdong Fenglu Aluminum Co., Ltd, Foshan 528133, China

M AN U

b

SC

China

ABSTRACT 1

Microstructural and property evolution of the Al-Zn-Mg-0.10%Sc-0.10%Zr alloy sheet

during its preparation were investigated in detail by means of optical microscopy (OM), scanning electron microscope (SEM), energy dispersive X-ray (EDX), transmission electron microscopy

TE D

(TEM), Vickers micro-hardness test and room temperature tensile test. Stress corrosion cracking (SCC) behavior of the Al-Zn-Mg-0.10%Sc-0.10%Zr alloy under different heat treatments was studied using slow strain rate test. The results showed that serious dendritic segregation existed in as-cast condition. The suitable homogenization treatment for Al-Zn-Mg-0.10%Sc-0.10%Zr alloy

EP

was 470°C/24h. After homogenization treatment, dissoluble Zn and Mg enriched non-equilibrium

AC C

phases dissolved into α-Al matrix completely. The suitable solid solution-aging treatment for AlZn-Mg-0.10%Sc-0.10%Zr alloy was solution treated at 470 °C for 60 min, followed by water quenching and then aged at 120 °C for 24 h. Under this aging temper, the grain structures were composed of sub-grains, η′ phases and nanometer-sized, spherical Al3(Sc, Zr) particles. Grain boundary precipitates (GBPs) area fraction was found to be an important parameter to evaluate the SCC susceptibility. The improved corrosion resistance from increasing aging temperature or prolonging aging time was due to the discontinuous η precipitates along the grain boundary and

* Corresponding author. Tel. /fax: +86 731 88830933. E-mail address: [email protected]. 1

ACCEPTED MANUSCRIPT the high area fraction of GBPs. The main strengthening mechanisms of Al-Zn-Mg-0.10%Sc0.10%Zr alloy are precipitation strengthening derived from η′ precipitates, dispersion strengthening, sub-grain strengthening and grain refinement caused by coherent Al3(Sc, Zr) particles.

RI PT

Key words: Al-Zn-Mg-Sc-Zr alloy; Al3(Sc, Zr) particle; microstructure; mechanical properties; stress corrosion cracking 1. Introduction

SC

Al-Zn-Mg alloys are considered as high strength Al alloys which are widely used for structure components in aerospace field, automobile manufacturing and advanced weapons

M AN U

systems because of their high strength-weight ratio, high fracture toughness and resistance to fatigue [1-4]. These alloys have a tensile strength above the strongest of the 6xxx series alloys and attain the highest strength of all Al alloys, exceeding normal structural steel. However, these alloys generally have poor ductility and low fracture strength in the as-cast condition, and extensive processing, which includes a combination of heat treatment and hot-cold working, is

TE D

required to improve the mechanical properties [5]. Before the 1960s, peak-aging treatment was used for Al-Zn-Mg alloys in order to maximize their strength. However, the precipitates at the grain boundaries after peak-aging treatment tended to form a continuous chain, resulting in high

EP

susceptibility to stress corrosion cracking (SCC) and relatively low fatigue crack growth resistance. Various heat treatments such as T76, T736 (T74), T7351, and retrogression and reaging (RRA)

AC C

have been extensively studied and developed to improve the strength, SCC resistance, and fatigue crack growth resistance of the Al-Zn-Mg alloys. Of late, material workers have found an effective measurement to modify microstructures and improve properties of Al alloys by microalloying with transition metals or rare earth metals (such as additions of Sc, Zr, Yb, Cr and Er) [6-7]. On a per-atom basis, Sc and Zr offer the greatest benefits, such as excellent grain refinement and the inhibition of recrystallization, mainly arising from the formation of extremely fine, coherent Al3(Sc, Zr) particles with L12 structure [8-12]. There are three ways that the Al3(Sc, Zr) particles can form in the Al alloy generally. (i) Primary Al3(Sc, Zr) particles that formed during solidification. They act as heterogeneous nucleation sites for α-Al and refine the as-cast grain

2

ACCEPTED MANUSCRIPT structure. (ii) Secondary Al3(Sc, Zr) dispersoids formed during homogenization and thermal mechanical process. These particles play a high anti-recrystallized effectiveness in the alloy [13]. (iii) Coherent Al3(Sc, Zr) phase precipitated during aging process. The precipitation temperature for Al3(Sc, Zr) ranges from 250 ºC to 350 ºC [14].

RI PT

The reports about the aluminum alloys with Sc and Zr additives are available but mainly concentrated on the precipitation behaviors of Al3(Sc, Zr) particles [15-16]. Desirable service performance from additions of minor Sc and Zr in aluminum alloys is mainly attributed to the presence of coherent, L12-ordered Al3(Sc, Zr) particles [17-18]. Jeoung et al. [19] observed that

SC

addition of 0.1 wt.% Sc into AA7075 was not enough to provide sufficient nucleation sites for dynamic recrystallization; however, the addition of 0.30 wt.% Sc resulted in significant grain

M AN U

refinement and fully equiaxed microstructure. M. Schöbel et al. [20] showed that Al alloys with additions of Sc and/or Zr exhibited a reasonably stable grain structure due to a uniform distribution of coherent Al3(Sc, Zr) precipitates that forms at temperatures above 300 ºC. N.Q. Tuan et al. [21] researched the microstructure and age-hardening behavior of Al-Sc alloy. At the aging temperature of 325 ºC, the average diameter is 4.3 nm for Al3Sc precipitates at the aging

TE D

peak. Under the temperature of 350 ºC and 7 days aging, the average diameter of Al3Sc precipitates is 13.7 nm. Sauvage et al. [22] observed precipitate stability and recrystallisation in the weld nuggets of friction stir welded Al-Sc alloys, found that Al3Sc precipitate size and density

EP

are unchanged in the nugget comparing to the base metal. These precipitates strongly reduce the boundary mobility of recrystallised grains, leading to a grain size in the nugget much smaller.

AC C

Three new Al-Zn-Mg-Sc-Zr alloys in the high-strength thermally strengthened weldable alloys based on the Al-Zn-Mg system with Sc and Zr additives are developed, namely, 01970, 01975 and 01981 [23-24]. However, the problem is that the additions of Zr and, especially, Sc are expensive and the solubility of these two elements in aluminum is rather low. Fortunately, the chemical composition of 01975 alloy and 01970 alloy is very close, the only difference is the Sc content in 01975 alloy is low, just about 0.07%. Furthermore, the low Sc alloy has good ductility and quenching process can be carried out after extrusion immediately. Low Sc alloy can achieve high strength and toughness, high stress corrosion resistance, excellent weld-ability. It has been developed into a new aerospace lightweight high strength structural material. However, up to now, 3

ACCEPTED MANUSCRIPT little research work has been devoted to the Al-Zn-Mg alloy with trace Sc addition. For its extensive applications, it is highly significant to figure out the evolution of microstructure and properties during preparation process of the alloy plate. The aim of this study is to obtain the AlZn-Mg-0.10%Sc-0.10%Zr alloy sheet with good mechanical properties. Therefore, the

RI PT

optimization of homogenizing treatment and solid solution aging treatment were investigated. Based on this, the evolution of microstructure and properties, the main mechanism and corrosion resistance behavior were also studied. 2. Experimental procedures

SC

The aluminum alloy used in the present work was Al-Zn-Mg-0.10%Sc-0.10%Zr alloy. Ingot with dimensions of 260 mm × 150 mm × 30 mm was made by the use of commercial pure metals

M AN U

(Al, Zn and Mg) and Al-Cu, Al-Mn, Al-Sc, Al-Zr master alloys, via semi-continuous casting in a crucible furnace, the chemical composition of the investigated alloy is shown in Table 1. Homogenization temperature was determined preliminarily on the basis of the upper limit of homogenization temperature obtained from differential scanning calorimetric (DSC) analysis. Specimens with dimensions of 20 mm × 20 mm × 10 mm were wire-cut from the center of ingot.

TE D

The specimens were homogenized at 460 °C, 465 °C, 470 °C, 475 °C and 480 °C for 24 h, then homogenized for 8, 16, 24 and 32 h at the optimized temperature, respectively. All the homogenized specimens were air-cooled to room temperature and investigated by microstructural

EP

observations. According to the microstructures and DSC analysis, optimized homogenization parameters were obtained. After optimal homogenization treatment, the ingots were annealed at

AC C

450 °C for 2~3 h, then hot and cold rolled to produce 2.5 mm thick plates immediately, with a total deformation of 92%. The cold-rolled sheet was subjected to solution treatment for 20 min ~ 2 h at 460 ºC 470 ºC 480 ºC 100 ºC

490 ºC, followed by water quenching, then aged for 4 ~ 48 h at

120 ºC 140 ºC 160 ºC. The detailed conditions of solution-aging treatments are listed

in Table 2. The solution-aging treated specimens under different conditions were investigated by tensile tests and microstructural observations. According to the microstructure and properties, optimized solution-aging process was acquired.

4

ACCEPTED MANUSCRIPT DSC analyses were conducted using a METZSCH DSC 200F3 differential scanning calorimeter. Samples with dimension of Φ5.5 mm × 4 mm were heated in an inert flowing atmosphere (N2) at a constant heating rate of 10 ºC/min from 25 to 650 ºC. The mechanical testing utilized transverse orientation (normal to rolling direction) specimens. The tensile tests were

RI PT

conducted on MTS810 tester with a gauge length of 28 mm. Yield strength represents 0.2% proof strength value as computed by the computer program controlling the machine. Three tensile tests were performed in each alloy and average values were used for discussion. Hardness tests were finished on a 401MVDTM digital Vickers micro-hardness with a load of 1960 mN and a dwell time

SC

of 10 s. Slow strain rate test was carried out to study the stress corrosion cracking (SCC) resistance according to GB15970.7-2000 [25].

M AN U

Detailed microstructure was carried out using an optical microscope (OM), scanning electron microscope (SEM) and transmission electron microscopy (TEM). Metallographic specimens were electrolytically polished and then anodized in the Baker’s solution (5 ml HBF4 + 217 ml H2O) at a best electroetching electric potential 22 V for 1~3 min. SEM samples were observed on a Quanta MK2-200 SEM, operating at 20 kV. Foils for TEM were cut from plates to characterize the

TE D

distribution of precipitates and substructures, subsequently ground to less than 100 µm and punched into 3 mm discs, and prepared by twin-jet electro-polishing with an electrolyte solution consisting of 25% HNO3 and 75% methanol (vol. %) below -25 ºC. TEM observations were

3. Results

EP

carried out on a TECNAI G220 electron microscope with a 200 kV accelerating voltage.

AC C

3.1. Microstructure of cast alloy

Optical micrograph of the grain structures of the alloy is presented in Fig. 1(a). The alloy

under as-cast condition consists of typical large, dendritic α-Al phase. Fig. 1(b) shows intermetallic phases at the grain boundaries. It can be seen serious dendritic segregation exists. EDX analysis results in Table 3 reveal that the white phases in point A are close to T-Mg32(Al, Zn)49 in composition, only Al, Zn and Mg elements are detected by EDX at triangular grain boundary area, suggesting that it may be the mixture of α-Al, η-MgZn2 and T-Mg32(Al, Zn)49. The grey phases as shown in point B are indissoluble impurity phases containing Fe, Mn and Si elements and they may be the α-Al matrix with solute of elements Zn, Mg, Fe, Si [26-28]. The concentration of the 5

ACCEPTED MANUSCRIPT elements in the alloy decreases from grain boundary to inside. Besides, massive intermediate compounds affect the subsequent processing properties. Therefore, a homogenization treatment is necessary to eliminate severe dendritic segregation in as-cast alloy. Higher temperature makes it easier to eliminate the segregation. Meanwhile, it increases the possibility of burnt phenomenon,

RI PT

so it is very important for practical application to obtain homogenization temperature. Fig. 2 shows the DSC analysis result of the alloy ingot. There is a clear endothermic peak at temperature 476.7 ºC. It has been discussed by Liu et al. [29] as well as by He et al. [30] that the peak corresponds to the melting of α-Al matrix and η-MgZn2, which implies that the start melting

SC

temperature (Tm) of the alloy ingot is 476.7 ºC. Considering the accuracy of furnace, the temperatures ranging from 460 ºC to 480 ºC were selected in this paper.

M AN U

3.2. Homogenization treatment

The purpose of homogenization treatment is to dissolve the non-equilibrium eutectic phases in α-Al matrix as much as possible, in order to improve the plastic workability, toughness and fatigue life of alloys. The distribution of intergranular phases in the ingot homogenized for 24 h by 460 °C ~ 480 °C are observed, as shown in Fig. 3. With increasing the homogenization

TE D

temperature, the volume fraction of non-equilibrium phases significantly decreases, the constituent phases at grain boundaries become discontinuous, showing a necklace distribution and the grain boundaries become thinner and clearer. When homogenized at 470 °C, the size of intergranular

EP

phases decrease and the intergranular phases distribute discontinuously. When the temperature is increased to 480 °C, the melting compounds both at grain boundaries and in triple conjunctions

AC C

can be observed, and the specimen is over-burnt. Fig. 4 shows the distribution of intergranular phases in the ingot homogenized at 470 °C for 8

h ~ 32 h. With prolonging of homogenization time, the dendritic-network structure reduces in the volume fraction, the constituent phases evolve to be discontinuous particles and the massive residual phases become smaller and sparse. When the specimen is homogenized for 24 h, the secondary phases are almost dissolved into the α-Al matrix (Fig. 4 (c)). Then, its amount decreases slightly with the further prolonging of homogenization time. Fig. 5 gives DSC curves of alloy at different homogenization temperature and time. The endothermic peak at 476.7 °C disappears gradually with increasing the homogenization temperature and prolonging of homogenization time. 6

ACCEPTED MANUSCRIPT It is corresponding to the dissolution of some non-equilibrium phases during homogenization, which is consistent with the facts in Fig. 3 and Fig. 4. As a result, the suitable homogenization treatment for Al-Zn-Mg-0.10Sc-0.10Zr alloy is 470°C/24h. 3.3. Solution-aging treatment

RI PT

Fig. 6 and Fig. 7 show the effect of solution treatment on microstructure and mechanical properties of the cold-rolled sheet aged at 120°C/24h. It is known that a higher isothermal temperature promotes a faster dissolution and an increased solute concentration in the matrix. On the other hand, solution treated temperature has to be limited below the eutectic temperature to

SC

avoid over-burnt. Typical fibrous structure exists along rolling direction after cold-rolled and massive non-equilibrium phases formed during rolling process (Fig. 6 (a)). With increasing

M AN U

solution temperature (Fig. 6 (b~e)) and prolonging homogenization time (Fig. 6 (f~i)), the volume fraction of the second phase in α-Al matrix reduces gradually. After solution treatment at 470 °C for 60 min, the alloy still keeps fibrous rolling deformation organization and the particles dissolve into matrix completely, only little impurity phases remain. However, after solution treatment at 490 °C for 60 min or 470° C for 90 min ~ 120 min, fibrous rolling deformation organization

TE D

abates and partial recrystallization occurs. With increasing solution temperature, strength and elongation of aged alloys increase firstly and then decrease (Fig. 7 (a)). As further prolonging holding time at 470 °C, strength increases firstly and then decreases, and elongation increases

EP

monotonously (Fig. 7 (b)). Fig. 8 shows the microstructure of solution treated sheet. After solution treatment at 470 °C for 60 min, unrecrystallized and fiber-like structures which are made up of

AC C

micrometer-sized sub-grains still remain in the alloy. Observed through high magnification, numerous, fine, nanometer-sized, spherical Al3(Sc, Zr) particles distribute in sub-grain and subgrain boundaries. Those particles strongly pin sub-grain boundaries, leading to structures maintaining unrecrystallization after rolling and solution treatment. Based on the above results, the best solution treatment for Al-Zn-Mg-0.10%Sc-0.10%Zr alloy is 470°C/60min. The effect of aging treatment on tensile properties of the sheet after solution treatment at 470 °C for 60 min is shown in Fig. 9. The tensile properties of the alloy aged at different temperatures for 24 h are shown in Fig. 9 (a). With increasing aging temperatures, strength rises firstly and then decreases, but elongation decreases monotonously. Fig. 9 (b) shows tensile 7

ACCEPTED MANUSCRIPT properties of the alloy aged at 120 °C, as a function of aging time. It can be seen that the age strengthening effect is obvious. In initial stage of aging, the strength rises rapidly. Peak aging is achieved at 24 h of aging. During the over-aging stage, the strength has a little change for the duration of aging time, but the elongation decreases monotonously. Comparatively, the sheet aged

RI PT

at 120 °C for 24 h exhibits better tensile properties. After suitable solution treatment (470°C/60min) and proper aging treatment (120°C/24h), the tensile strength (MPa), yield strength (MPa) and elongation (%) of the sheet reach 555 ± 2, 524 ± 4 and 12.3 ± 0.6 respectively.

The microstructure evolution of aged alloy is observed by TEM. Fig. 10 shows TEM

SC

micrographs of the sheet aged at different temperatures for 24 h and 120 °C for different durations. Nanometer-sized, spherical Al3(Sc, Zr) particles can be observed in grain after aging treatment. In

M AN U

addition, a large number of fine precipitates distribute homogeneously in α-Al matrix (Fig. 10 (c)). For aging at 100°C/24h (Fig. 10 (a)) or 120°C/4h (Fig. 10 (e)), high density of precipitates are observed in both conditions. Under those aging conditions, GP zones are the main strengthening precipitates in the alloy [31-33]. According to Refs. [34-37], sizes of GP zones in those conditions rang from 2 to 8 nm and they are coherent with α-Al matrix, with internal ordering of Zn and

TE D

Al/Mg on the {200}Al planes. With increasing aging temperature and prolonging aging time (aging at 120°C for 12~24h), GP zones gradually transfer into a mass of dispersed distribution η′ phase; these precipitates grow, aggregate and accelerate along grain boundaries during the aging process,

EP

as demonstrated in Fig 10 (b, g). Meanwhile, the nucleation/growth of GP zones and η′ phase becomes much more apparent, which represents a significant strengthening effect. After 24h of

AC C

aging, the microstructures are still constituent with micro-scaled sub-grains. Besides, short rodlike particles can be seen in the α-Al matrix, and some coarsen rod precipitates distribute discontinuously along grain boundaries (Fig. 10 (b)). As is well known, η precipitates are the main grain boundary precipitates in Al-Zn-Mg alloys. After aging at 120°C for 48h, more precipitates transform from the metastable η′-(MgZn2) phases to the stable η phases. The η′ phases can be observed within the grain, the η-MgZn2 precipitates coarsening and completely discontinuous along the grain boundary, as demonstrated in Fig. 10 (i). However, the stable η phases are incoherent with the α-Al matrix and cannot improve the tensile strength. Meanwhile, the distributed spacing and grain size of boundary precipitates become larger with increasing aging 8

ACCEPTED MANUSCRIPT time, as demonstrated in Fig. 10 (c, d), leading to the decrease of strength in sheets as shown in Fig. 9 (b). So PA alloy has the highest Rm and Rp, while the Rm and Rp values of OA alloy decreases (Fig. 9). Additionally, the width of PFZ after (140°C/24h) and (120°C/48h) aging treatments increases to be 40.4 nm and 48.7 nm, respectively. Generally, aluminum alloy with this

RI PT

type of grain boundary structure demonstrates higher corrosion resistance property. The SAD pattern of the alloy at 120°C/24 h in [011]Al projection is shown in Fig. 10 (f), the diffraction spot from α-Al matrix has been indexed, these spots are Al3(Sc, Zr) particle with a Ll2 cubic crystal structure, which is coherent with α-Al matrix. These Al3(Sc, Zr) precipitates, which range in size

SC

from 20 to 30 nm, help in stabilizing the fine grained microstructure and strengthening the alloys by pinning the grain boundary and hindering the dislocation motion [38-39].

M AN U

The evolution of the Al3(Sc, Zr) precipitates can be divided into three stages: nucleation, growth and coarsening. At the first stage, clusters of Al and Sc atoms are formed to a critical radius (r*) for nucleation. The nucleation rate is given by [40]:
= 

 

 [

∗ ∗ 

]



(1)

where h and k are the Planck and Boltzmann constants, Q* is the activation energy, σ is surface

TE D

energy, Nv is the number of nucleation in unit volume, respectively. The nucleation rate heavily depends on T. An increase in temperature leads to a decrease in nucleation rate as well as number density if temperature is above a critical value. At the second stage, initial Al3Sc precipitates starts

EP

to grow. The growth is controlled by the diffusion of Sc and Zr atoms. Sc atoms diffuse greatly faster than Zr atoms do. The growth is due to the absorption of rapid diffusing Sc firstly. After half

AC C

of the Sc is removed into the precipitates, the Zr depletion starts to increase [41], resulting in a Zrriched shell on Al3Sc precipitates. At the third stage, small precipitates dissolve and then condensate their atoms to bigger ones. Slow diffusing Zr atoms in the Al3(Sc, Zr) precipitates shell have to be dissolved first before the removal of the whole precipitates. According to LSW (Lifshitz, Slyozov and Wagnar) theory, the coarsening behavior of Al3(Sc, Zr) precipitate is illustrated by Eq. (2) [42]:  = 

   &'  #$% #  &' )  $% ][ ) ]} ! {[ ($% ()

(2)

9

ACCEPTED MANUSCRIPT where αkv is the coarsening constant, Vm is the molar volume of one Al3(Sc, Zr) precipitate, σ is an isotropic interfacial free energy, R is the gas constant, T is the absolute temperature, DSc and DZr are the diffusion coefficients of Sc and Zr, respectively, #+ is i-th component in the matrix, ki is the distribution coefficient of the i-th element between the matrix and Al3(Sc, Zr) phase. As is

RI PT

known to all, the low interface energy and strain energy between the dispersion phase of L12 type structure and the α-Al matrix can reduce the coarsening rate of dispersion phases. From the TEM image, a large number of small, dispersed L12 type Al3(Sc, Zr) particles can be seen and the parameters between Al3(Sc, Zr) and α-Al matrix is similar. So according to the above theory,

which is in agreement with the conclusion we got before.

SC

Al3(Sc, Zr) particles have high thermal stability and are not easy to grow at high temperature,

M AN U

However, one TEM image of the grain boundaries cannot present all characterization of grain boundaries accurately. Even for the same alloy under same temper, the size of grain boundary precipitates (GBPs), the width of PFZ, area fraction of GBPs and the size of matrix precipitation phases (MPTs) all show differences more or less from one area to another. In order to obtain the accurate information of the grain boundary microstructures, more than 80 GBPs in different grain

TE D

boundary areas are measured by software and analyzed by data statistics, the results are listed in Table 4. With the increasing of aging temperature, the PFZ width and the average size of GBPs all increase. The value of area fraction of GBPs as an aggregative index is complexly influenced by

EP

the size and number of GBPs, grain boundary width and the distributed spacing of GBPs. The area fraction of GBPs increases from 11.3% to 23.3%. And the MPTs size increases from 2.5 nm to

AC C

11.2 nm, as well. Moreover, the average values of GBPs size, PFZ width and GBPs area fraction increase with the increasing of aging time. The MPTs size also increases significantly. 3.4. Age-hardening and the conductivity variation Fig. 11 shows the aging hardening behavior of the alloy aged at 120 °C for varying time. The

variation of hardness with aging time displays the characteristic age harden behavior of aluminum alloys, i.e. progressive increase of hardness with aging time, reaching to maximum value (peak aging condition) and then decrease of hardness value. As seen in Fig. 11, it can be easily found that the hardness of the alloy dramatically increases during the initial stage of aging processing, which is attributed to the formation of the large numbers of GP zones and metastable η′ 10

ACCEPTED MANUSCRIPT precipitates. The aging curve displays that the peak hardness has been attained at about 24 h which is attributed to the precipitation of large numbers of fine η′ precipitates in the α-Al matrix, as shown in Table 4. With increasing aging time, the hardness of the alloy stays gentle for a relatively long time and then decreases, the decreases at longer aging time is due to gradual

RI PT

transformation of the metastable η′ phases to the stable η phases and subsequent growth of η′ precipitates. The conductivity of the alloy increases persistently almost all through the aging process and a peak value of 35.3 %IACS is observed at 36 h, as shown in Fig. 11. Thereafter, a slight descent (around 0.7 %IACS) is observed at the terminal stage. The long-lasting increase in

SC

the conductivity is quite similar to that observed in hardness case, the slight decrease at the terminal stage, however is somewhat unusual. The SCC resistance has been found improved with

M AN U

the increase of electrical conductivity in varied Al-Zn-Mg alloys, particularly from the peak-aged to the over-aged condition [43-44]. The SCC resistance of specimens in the current study is preliminarily evaluated based on the conductivity. The conductivity of the peak-aged condition is about 27.48 %IACS, much lower than the 34.61 %IACS of 120°C/48h, suggesting a higher SCC resistance, as compared with that obtained in 120°C/24h condition, can be obtained through

3.5. SCC resistance

TE D

prolonging aging time.

Fig. 12 shows the typical SSRT nominal stress and nominal curves of the alloy tested in air

EP

and in aerated 3.5 wt.% NaCl solution. With the increase of the strain, the stress increases rapidly and then holds constant or decreases to some extent after reaching the peak stress value. Through

AC C

the SSRT, the susceptibility index of the stress corrosion of the studied alloys is derived. According to the obtained results and the literatures [45-46], Eq. (3) is used to evaluate the index of susceptibility to SCC.

RSCC = (1 −

Rsol ) × 100% Rair

(3)

where Rsol and Rair mean the values of any of each measured property in air and corrosive solution, respectively, three measured properties were used: yield strength (Rp), ultimate tensile strength (Rm) and tensile elongation (A). Rscc means the SCC susceptibility evaluated by different property parameters. The value of Rscc decreases, indicating that the SCC resistance increases. Under the 11

ACCEPTED MANUSCRIPT same aging condition, compared with the elongation of the sample tested in air, the elongation of the sample tested in 3.5 wt.% NaCl solution is lower. Besides, with increasing aging temperature and prolonging aging time, the elongation increases. According to Eq. (3), the SCC susceptibilities of the studied alloy under different aging

RI PT

tempers are calculated and the results are listed in Table 5 and Fig. 13. It is apparent that the SCC susceptibilities of the investigated alloy under different aging conditions are significantly different. It can be found that increasing aging temperature or prolonging aging time, the values of Rscc decrease significantly, which indicates that SCC resistance can be improved dramatically by

SC

increasing aging temperature or prolonging aging time. As shown in Fig. 13, the UAII alloy has the highest elongation loss and lowest SCC resistance. The beneficial effect of over-ageing on

M AN U

SCC resistance is observed for OAI~OAIV alloys, which almost have no SCC susceptibility in aerated 3.5 wt.% NaCl solution. Y. Reda [47] observed that, the SCC resistance of tradition OA condition Al-Zn-Mg alloys can be improved at the expense of a 10-20% strength loss when compared with the PA condition alloy. However, in the present work shown in Table 5 for Al-ZnMg-0.10Sc-0.10Zr alloy, the SCC resistance is significantly improved but accompanied by the Rp)

2%). In summary, the corrosion susceptibility of Al-Zn-Mg

TE D

strength loss slightly (Rscc(Rm,

alloy decreases without sacrificing much strength by adding Sc and Zr and adequately increasing aging temperature or prolonging aging time.

EP

SEM images of the fracture surface of the studied alloy under different aging tempers after SSRT in air and 3.5 wt.% NaCl solution are shown in Fig. 14 and Fig. 15. All the materials have

AC C

undergone severe plastic deformation during rolling process and spherical Al3(Sc, Zr) strongly inhibits recrystallization. Therefore, it represents strong anisotropy, a preferred dimples or voids growth can be seen along the horizontal direction of samples. It also shows that the drop in ductility is associated with a transition in fracture mode. The fracture surface of the specimens tested in air shows only dimples characteristic of ductile fracture, as illustrated in Fig. 14 (a, c) and Fig. 15 (a). In contrast, large area fractions of brittle intergranular fracture and some small areas with cleavage-like fracture are present on the fracture surfaces of the specimens tested in 3.5 wt.% NaCl solution. The stress corrosion cracks can be seen along the grain boundaries of samples tested in 3.5 wt.% NaCl solution (Fig. 14 (b)). Microvoids act as stress concentrators, and these 12

ACCEPTED MANUSCRIPT sites can lead to the initiation of cracks. They grow until adjacent voids connect, and coalesce into large voids until final failure of the material under the stress, which explains why the UA alloy is very susceptible to corrosive environment. With the increase of aging degree, the equiaxial dimples become more and deeper, as shown in Fig. 15 (c, e, f), the OA state alloy hard to be

RI PT

corroded to failure, and the SCC resistance is improved, which is consistent well with the results in Table 5. 4. Discussion

4.1. Evolution of microstructures and properties during homogenization and solution-aging

SC

treatment

Owing to the difference of cooling rate between the exterior and center of ingot, lots of Zn

M AN U

and Mg enriched nonequilibrium phases and indissoluble impurity phases containing Fe and Si elements form, therefore, serious dendritic segregation exists in the as-cast alloy, as illustrated in Fig. 1 (b). A necessary homogenization treatment is designed to remove micro-segregation and dissolve large soluble non-equilibrium eutectic phases. In addition, because of the investigated alloy contains trace amount of Sc and Zr, obtain massive Al3(Sc, Zr) particles which can prevent

TE D

re-crystallization during hot rolling and solution treatment is the another purpose of homogenization heat treatment. On the base of DSC analysis and distribution of intergranular phases, the suitable homogenization temperature for Al-Zn-Mg-0.10%Sc-0.10%Zr alloy is 470 °C.

EP

With increasing the homogenization temperature from 460 °C to 475 °C and prolonging homogenization time from 8 h to 24 h, the volume fraction of non-equilibrium phase decreases

AC C

gradually, the constituent phases at grain boundaries become discontinuous showing a necklace distribution and the grain boundaries become thinner and clearer. After 24 h of homogenizing at 470 °C, dissoluble nonequilibrium phases are primarily dissolved into α-Al matrix and the targets of homogenization treatment are achieved. When the temperature is increased to 480 °C, the specimen is over-burnt. As further prolonging homogenization time to 32 h, the amount of nonequilibrium phase decreases slightly. As a result, the suitable homogenization treatment is 470°C/24h. In order to obtain a higher strength or a larger driving force for second phase precipitation during aging treatment of aged alloy sheets, a larger degree of super-saturation must be gotten. It 13

ACCEPTED MANUSCRIPT means the coarse phases formed during the cooling process should be dissolved into the α-Al matrix as much as possible. Massive, fine, nanometer-sized Al3(Sc, Zr) particles strongly pin subgrain boundaries and effectively prevent re-crystallization during hot and cold rolling process and heat treatment, leading micro-scaled sub-grains structures remain in aged sheets, which results in

RI PT

typical fibrous structure (Fig. 6) still can be found after solution-aged and the strength of Al-ZnMg-0.10%Sc-0.10%Zr alloy is 40 MPa higher than normal Al-Zn-Mg alloy [48]. When soluted at lower temperatures (460 °C) or holding for shorter time (20 min ~ 40 min), the aged sheets are obtained lower strength (Fig. 7) for the reason that lots of phases still exist in the alloy and the

SC

plates are gained lower degree of solid solubility under such solution condition. Phases dissolve into the α-Al matrix at higher temperatures (490 °C) or holding for longer time (90 min ~ 120

M AN U

min). However, fibrous rolling deformation organization abates and partial recrystallization occurs after solution, also result in the lower strength of aged alloy. As a whole, strength of the studied alloy increases firstly and then decreases. Based on the above results, the best solution treatment is 470°C/60min.

The strengthening of the alloy is usually considered the summation resulting from different

σtot = σ0 + σss + ∆σprec

TE D

contributions:

(4)

where σ0 is the friction stress of the α-Al matrix and is considered constant, ∆σprec is the

EP

precipitates contribution, while σss is the solid solution contribution. If we assume that the contribution from each element is additive, the solution hardening potential can be calculated as

AC C

follows [49]: ∆σss = ΣKiXi2/3

(5)

where Xi is the concentration of a specific alloying element in the matrix and Ki is the corresponding scaling factor. The main precipitation sequence of Al-Zn-Mg alloys is usually summarized as follows: supersaturated solid solution (SSS)→GP zones→metastable η′→stable η. For Al-Zn-Mg alloys, it is common to apply two step artificial aging treatments to achieving high strength while maintaining good stress corrosion cracking resistance. The compositions of the binary phases η′ and η are dependent on heat treatment and alloy composition. The GP zones and metastable η′ phase are the main aging-hardening precipitates in commercial Al-Zn-Mg alloys. 14

ACCEPTED MANUSCRIPT The aging peak in Fig. 11 is the result from the high-density GP zones (especially GP II zones) and η′ phases [50]. As the degree of aging treatment increasing, the precipitates transform gradually from the metastable η′ phase to the stable η phase. Unfortunately, the η phase is equilibrium phase, non-coherent with α-Al matrix and does not contribute strongly to the

RI PT

strengthening effect of the alloy, so the strength reduces at OA states. In addition, Al3(Sc, Zr) particles do not change the characteristic of aging precipitation during aging treatment in the studied alloy. As is well known, Al3(Sc, Zr) particles are coherent with α-Al matrix, the strengthening caused by Al3(Sc, Zr) particles and other precipitates can be estimated by Orowan

SC

bowing mechanism. The increase of yield strength, Rp, can be calculated by the Ashby-Orowan equation [51]: <

M AN U

/.123

-. = 4(& 6)&/ 9 :; 3

(6)

where M is the Taylor factor, M = 3.06; G is shear modulus of the α-Al matrix; G = 27.8 GPa; b is the Burgers vector of dislocation in Al, b = 0. 286 nm; ν is the Poisson ratio, for Al, ν = 0.331, and λ is the interspacing of particles. 4

(7)

TE D

λ = r ( = )&/

where r is the radius of particles and f is the volume fraction of particles. The average grain size of Al3(Sc, Zr) particles is about 20-30 nm; the volume fraction of Al3(Sc, Zr) particles is about 4.5 ×

EP

10-3 in the studied alloy. Therefore, the contribution of the dispersoids to the yield strength (Rp) can be calculated by using the measured parameters of dispersoids in the material [5, 52]. The

AC C

calculated result is about 30 MPa.

Actually, the grain boundary strengthening is also an important strengthening mechanism for

Al-Zn-Mg-0.10Sc-0.10Zr alloy. Based on the above discussion, Al3(Sc, Zr) particles can effectively refine the grain size. After PA treatment, the grain sizes are 2.95 µm in Al-Zn-Mg0.10Sc-0.10Zr alloy. Combined with the research reported by us [48], the grain sizes are 9.58 µm in Al-Zn-Mg alloy. The standard Hall-Petch equation, Eq. (8), is employed to relate the yield strength of the material (σ) to the average grain size (d, µm). σ = σ0 + kd-1/2

(8)

15

ACCEPTED MANUSCRIPT where σ0 is the intrinsic resistance of the lattice to dislocation motion and k is a parameter that describes the relative strengthening contribution of grain boundaries, k = 68 MPa × µm1/2 [53-54]. The increase of yield strength can be calculated by Eq. (6), the calculated result is about 18 MPa in Al-Zn-Mg -0.10Sc-0.10Zr alloy. Therefore, the calculated YS increases from Sc and Zr

RI PT

microalloying additions under PA state were about 48 MPa, which approached the experimental results we got before (53MPa) [48]. From Fig. 10, the grain structures of PA alloy are composed of sub-grains, η′ phases and Al3(Sc, Zr) particles. In short, the main strengthening mechanisms of Al-Zn-Mg-0.10%Sc-0.10%Zr alloy are precipitation strengthening derived from η′ precipitates,

SC

dispersion strengthening, sub-grain strengthening and grain refinement caused by coherent Al3(Sc, Zr) particles.

M AN U

4.2. Corrosion resistance

The present results in Fig. 13 and Table 5 clearly demonstrate that ageing treatment has a significant effect on the susceptibility to SCC of the investigated alloy. OA temper can improve the SCC resistance as observed for OAI ~ OAIV alloy, which almost has no SCC susceptibility in aerated 3.5 wt.% NaCl solution (Rscc(Rm, Rp) 2%), which is consistent with those discoveries by

TE D

many authors [46, 55]. A large number of fine precipitates (GP zones) form homogeneously within grains and along grain boundaries in UA tempers alloy (Fig. 10 (a, e, g)). With increasing aging temperature or prolonging aging time, short rod-like aging precipitates gradually coarsen in

EP

the α-Al matrix or along grain boundary. Under OA temper, PFZ forms, the spacing of grain boundary precipitates becomes larger and discontinuous.

AC C

As is well known, the corrosion resistance of Al-Zn-Mg alloys is associated with the distribution of grain boundary precipitates. Up to the present, GBPs, MPTs, PFZ and the area fraction of the GBPs have been discussed concerning the susceptibility of SCC. As literatures [5657] reported, the loss factor of elongation is generally chosen to describe the SCC susceptibility of samples, which is also used in this paper. Fig. 16 represents the relationship between SCC susceptibility and area fraction of GBPs of the studied alloy. Tsai and Chuang [58] reported that the size of matrix precipitates was the major factor affecting the SCC resistance when GBPs were larger than a critical size that could nucleate hydrogen bubbles, and OA treatment could improve the SCC resistance of Al-Zn-Mg alloys with the increase of MPTs size. The increasing size of 16

ACCEPTED MANUSCRIPT GBPs during OA temper has also been proposed to explain the lower SCC susceptibility in Al-ZnMg alloys. It can be seen that the studied alloy under OAII and OAIV treatment has the highest area fraction of GBPs and the lowest SCC susceptibility among all alloys with different heat treatments (Fig. 16). The results are consistent with results in Table 5 and Fig. 13.

RI PT

Two basic mechanisms have been proposed to model SCC of Al-Zn-Mg alloys: anodic dissolution and hydrogen embrittlement. Recently, it is reported that both anodic dissolution and hydrogen embrittlement operated in the SCC process. The precipitates in the grain boundaries are Mg-rich phases, which have the electrode potential different from the α-Al matrix. This results in

SC

the anodic dissolution and forming critical defects in the first stage of the SCC process in the 3.5 wt.% NaCl solutions. Furthermore, the hydrogen produces in the crack tip also leads to the

M AN U

hydrogen embrittlement in the grain boundaries. When precipitates distribute continuously at grain boundaries, shown in Fig. 10 (a, b, e), there exists a positive corrosion path resulting from the galvanic reaction between the anodic precipitates of precipitates at the grain boundaries and the αAl matrix at their adjacent periphery, which leads to its low resistance to corrosion. In the OA temper alloys whose area fraction of GBPs are sufficiently high (Fig. 10 (d, h, i)), the large size

TE D

and inter-particle spacing of the grain boundary particles decrease the anodic reaction rate. Furthermore, the large η phases in the grain boundaries can also act as the trapping sites for atomic hydrogen and create molecular hydrogen bubbles to reduce the concentration of the atomic

EP

hydrogen. Thus, the resistance of SCC is great under this temper. However, in PA temper, the precipitates in the grain boundaries are smaller than those in the OA states. That will increase the

AC C

anodic solution speed, reduce the number of the trapped hydrogen atoms, and increase the hydrogen atom diffusion on the grain boundary. Hence, the SCC resistance of the Al-Zn-Mg0.10%Sc-0.10%Zr alloy can be improved by increasing aging temperature or prolonging aging time.

5. Conclusions 1. Serious dendritic segregation exists in as-cast condition of Al-Zn-Mg-0.10%Sc-0.10%Zr alloy. A lot of Zn and Mg enriched non-equilibrium phases and indissoluble impurity phases containing Fe, Si, Mn elements are concentrated on grain boundaries.

17

ACCEPTED MANUSCRIPT 2. The suitable homogenization treatment for Al-Zn-Mg-0.10%Sc-0.10%Zr alloy is 470°C/24h. After homogenization treatment, non-equilibrium phases are dissolved into matrix and the dendritic-network structure disappears. 3. The suitable solution-aging treatment for Al-Zn-Mg-0.10%Sc-0.10%Zr alloy is solution treated

RI PT

at 470 °C for 60 min, followed by water quenching and then aged at 120 °C for 24 h. The grain structures of PA alloy are composed of sub-grains, η′ phases and nanometer-sized, spherical Al3(Sc, Zr) particles. The main strengthening mechanisms of Al-Zn-Mg-0.10%Sc-0.10%Zr alloy are precipitation strengthening derived from η′ precipitates, dispersion strengthening, sub-grain

SC

strengthening and grain refinement caused by coherent Al3(Sc, Zr) particles.

4. Grain boundary precipitate, matrix precipitation phase and precipitate free zone have been

M AN U

discussed as important parameters to evaluate the SCC susceptibility of alloys. Increasing aging temperature or prolonging aging time can enhance the SCC resistance of the Al-Zn-Mg-0.10%Sc0.10%Zr alloy, which is attributed to the discontinuous η precipitates along the grain boundary and the high area fraction of GBPs. Acknowledgements

Development

Program

TE D

The authors wish to acknowledge the financial supports of National Key Basic Research and of

China

under

Grant

(2012CB619503)

and

Science

and

References

EP

Technology Innovation Fund of Foshan City (2013HH100055).

AC C

[1] H.A. Godinho, A.L.R. Beletati, E.J. Giordano, C. Bolfarini, J. Alloys Compd. 586 (2014) 139142.

[2] Y.L. Deng, L. Wan, Y.Y. Zhang, X.M. Zhang, J. Alloys Compd. 509 (2011) 4636-4642. [3] N. Yazdian, F. Karimzadeh, M. Tavoosi, J. Alloys Compd. 493 (2010) 137-141. [4] S. Rajakumar, V. Balasubramanian, Mater. Manuf. Process. 27 (2012) 78-83. [5] O.N. Senkov, M.R. Shagiev, S.V. Senkova, D.B. Miracle, Acta Mater. 56 (2008) 3723-3738. [6] M.J. Jones, F.J. Humphreys, Acta Mater. 51 (2003) 2149-2159. [7] Y.L. Wu, F.H. Froes, C.G. Li, A. Alvearez, Metall. Mater. Trans. A 30 (1999) 1017-1024. [8] C. Booth-Morrison, D.C. Dunand, D.N. Seidman, Acta Mater. 59 (2011) 7029-7042.

18

ACCEPTED MANUSCRIPT [9] S. Dev, A.A. Stuart, R.C. Ravi Dev Kumaar, B.S. Murty, K.P. Rao, Mater. Sci. Eng. A 467 (2007) 132-138. [10] M.E. van Dalen, T. Gyger, D.C. Dunand, D.N. Seidman, Acta Mater. 59 (2011) 7615-7626. [11] A.K. Mukhopadhyay, A. Kumar, S. Raveendra, I. Samajdar, Scripta Mater. 64 (2011) 386-

RI PT

389. [12] A. Bommareddy, M.Z. Quadir, M. Ferry, J. Alloys Compd. 527 (2012) 145-151. [13] V. Zakharov, Met. Sci. Heat Treat. 53 (2012) 414-419. [14] J. Röyset, Metall. Sci. Technol. 25 (2007) 11-21.

SC

[15] M. Vlach, I. Stulíková, B. Smola, N. Zaludová, J. Cerná, J. Alloys Compd. 492 (2010) 143148.

M AN U

[16] N.A. Belov, A.N. Alabin, I.A. Matveeva, 583(2014) 206-213.

[17] C. Booth-Morrison, D.C. Dunand, Da.N. Seidman, Acta Mater. 59 (2011) 7029-7042. [18] N.Q. Vo, D.C. Dunand, D.N. Seidman, Acta Mater. 63(2014) 73-85. [19] J.H. Kim, J.H. Kim, J.T. Yeom, D.G. Lee, S.G. Lim, N.K. Park, J Mater Process Tech. 187188 (2007) 635-639.

TE D

[20] M. Schöbel, P. Pongratz, H.P. Degischer, Acta Mater. 60 (2012) 4247-4254. [21] N.Q. Tuan, A.M.P. Pinto, H. Puga, L.A. Rocha, J. Barbosa, Mater. Sci. Eng. A 601 (2014) 70-77.

EP

[22] X. Sauvage, A. Dédé, A. Cabello Muñoz, B. Huneau, Mater. Sci. Eng. A 491 (2008) 364-371, [23] Y. Deng, Z.M. Yin, J.Q. Duan, K. Zhao, B. Tang, Z.B. He, J. Alloys Compd. 517 (2012) 118-

AC C

126.

[24] V.G. Davydov, V.I. Elagin, V.V. Zakharov, T.D. Rostova, Met. Sci. Heat Treat. 38 (1996) 347-350.

[25] GB/T 15970.7-2000. National Standard of China. Corrosion of Metals and Alloys-Stress Corrosion Testing-Slow Strain Rate Testing. [26] P.F. Jia, Y.H. Cao, Y.D. Geng, L.Z. He, N. Xiao, J.Z. Cui, Mater. Sci. Eng. A 612 (2014) 335-342. [27] H.J. Wang, J. Xu, Y.L. Kang, M.G. Tang, Z.F. Zhang, J. Alloys Compd. 585 (2014) 19-24. [28] L.M. Wu, W.H. Wang, Y.F. Hsu, S. Trong, J. Alloys Compd. 456 (2008) 163-169. 19

ACCEPTED MANUSCRIPT [29] Y. Liu, D. Jiang, W.L. Xie, J. Hu, B. Ma, Mater. Charact. 93 (2014) 173-183. [30] L.Z. He, X.H. Li, P. Zhu, Y.H. Cao, Y.P. Guo, J.Z. Cui, Mater. Charact. 71 (2012) 19-23. [31] T. Hu, K. Ma, T.D. Topping, J.M. Schoenung, E.J. Lavernia. Acta Mater. 61 (2013) 21632178.

RI PT

[32] W. Guo, J.Y. Guo, J.D. Wang, M. Yang, H. Li, X.Y.Wen, J.W. Zhang, Mater. Sci. Eng. A 634 (2015) 167-175.

[33] T. Marlaud, A. Deschamps, F. Bley, W. Lefebvre, B. Baroux, Acta Mater. 58 (2010) 48144826.

SC

[34] P.K. Rout, M.M. Ghosh, K.S. Ghosh, Mater. Charact. 104 (2015) 49-60.

[35] L.K. Berg, J. Gjønnes, V. Hansen, X.Z. Li, M. Knutson-wedel, G. Waterloo, D. Wallenberg,

M AN U

Acta Mater. 49 (2001) 3443-3451.

[36] Z.H. Li, B.Q. Xiong, Y.G. Zhang, B.H. Zhu, F. Wang, H.W. Liu, Mater. Charact. 59 (2008) 278-282.

[37] C.B. Fuller, M.W. Mahoney, M. Calabrese, L. Micona, Mater. Sci. Eng. A 527 (2010) 22332240.

TE D

[38] A.K. Lohar, B. Mondal, D. Rafaja, Mater. Charact. 60 (2009) 1387-1394. [39] K.E. Knipling, R.A. Karnesky, C.P Le, D.C. Dunand, D.N. Seidman, Acta Mater. 58 (2010) 5184-5195.

EP

[40] Deschamps A, Lae L, Guyot P, Acta Mater. 55 (2007) 2775-2783. [41] J.D. Robson, Acta Mater. 52 (2004) 1409-1421.

AC C

[42] C.J. Kuehmann, P.W. Voorhees, Metall Mater Trans A, 27(1996) 937-938. [43] M.J. Starink, X.M. Li, Metall. Mater. Trans. A 34 (2003) 899-911. [44] J.T. Jiang, W.Q. Xiao, L. Yang, W.Z. Shao, S.J. Yuan, L. Zhen, Mater. Sci. Eng. A 605 (2014) 167-175.

[45] Y.J. Shi, Q.L. Pan, M.J. Li, X. Huang, B. Li, Mater. Sci. Eng. A 621 (2015) 173-181. [46] X.Y. Sun, B. Zhang, H.Q. Lin, Y. Zhou, L. Sun, J.Q. Wang, E.H. Han, W. Ke, Corros. Sci. 77 (2013) 103-112. [47] Y. Reda, R. Abdel-Karim, I. Elmahallawi, Mater. Sci. Eng. A 485 (2008) 468-475. [48] B. Li, Q.L. Pan, X. Huang, Z.M. Yin, Mater. Sci. Eng. A 616 (2014) 219-228. 20

ACCEPTED MANUSCRIPT [49] O.R Myhr, Ø Grong, S.J Andersen. Acta Mater. 49 (2001) 65-75. [50] G. Fribourg, Y. Bréchet, A. Deschamps, A. Simar, Acta Mater. 59 (2011) 3621-3635. [51] Y.J. Li, A.M.F. Muggerud, A. Olsen, T. Furu, Acta Mater, 60 (2012) 1004-1014. [52] T. Wang, T. Gao, P. Zhang, J.F. Nie, X.F. Liu, J. Alloys Compd. 589 (2014) 19-24.

RI PT

[53] N. Hansen, Scripta Mater. 51 (2004) 801-806. [54] P. Lehto, H. Remes, T. Saukkonen, H. Hänninen, J. Romanoff, Mater. Sci. Eng. A 592 (2014) 28-39.

[55] Y. Deng, Z.M. Yin, K. Zhao, J.Q. Duan, J. Hu, Z.B. He, Corros Sci. 65 (2012) 288-298.

SC

[56] X.Y. Sun, B. Zhang, H.Q. Lin, Y. Zhou, L. Sun, J.Q. Wang, E.H. Han, W. Ke, Corros Sci. 77 (2013) 103-112.

M AN U

[57] H. Lee, Y. Kim, Y. Jeong, S. Kim, Corros Sci. 55 (2012) 10-19.

AC C

EP

TE D

[58] T.C. Tsai, T.H. Chuang, Metall. Mater. Trans. A 27 (1996) 2617-2627.

21

ACCEPTED MANUSCRIPT Figure Captions: Fig. 1. Microstructures of as-cast alloy: (a) OM image; (b) SEM image. Fig. 2. DSC analysis of the Al-Zn-Mg-0.10%Sc-0.10%Zr alloy ingot. Fig. 3. Distribution of intergranular phases of different homogenization temperature for 24 h: (a)

RI PT

As-cast; (b) 460 °C; (c) 465 °C; (d) 470 °C; (e) 475 °C; (f) 480 °C. Fig. 4. Distribution of intergranular phases for different homogenization time at 470 °C: (a) 8 h; (b) 16 h; (c) 24 h; (d) 32 h.

Fig. 5. DSC curves of the alloy: (a) 24 h at different homogenization temperature; (b) 470 °C for

SC

different homogenization time.

Fig. 6. Microstructure of the studied alloys under different solution treatment: (a) cold rolling; (b)

M AN U

USI; (c) PS; (d) OSI; (e) OSII; (f) USII; (g) USIII; (h) OSIII; (i) OSIV.

Fig. 7. Mechanical properties of the alloy under different solution treatment (aged at 120°C/24h): (a) soluted for 60 min at 460 °C, 470 °C, 480 °C, 490 °C; (b) soluted at 470 °C for 20 min, 40 min, 60 min, 90 min, 120 min. Rm, ultimate tensile strength; Rp, yield strength; A, elongation.

TE D

Fig. 8. Microstructures of the alloy after solution treatment at 470 °C for 60 min: (a) Fiber-like structure with sub-grains; (b) Al3(Sc, Zr) within grains and in sub-grain boundaries; (c) SAD in [001]Al projection.

EP

Fig. 9. Mechanical properties of the alloy under different aging treatment after 470°C/60min solution treatment: (a) aged for 24 h at 100 °C, 120 °C, 140 °C, 160 °C; (b) aged at 120 °C

AC C

for 4 h, 12 h, 24 h, 32 h, 48 h. Rm, ultimate tensile strength; Rp, yield strength; A, elongation.

Fig. 10. Bright field microstructural evolution of the studied alloy under various aged conditions: (a) UAI; (b) PA; (c) OAI; (d) OAII; (e) UAII; (f) 120°C/24h, SAD in [011]Al projection; (g) UAIII; (h) OAIII; (i) OAIV. Fig. 11. Variation in the hardness and the relative electric conductivity of the alloy during aging process. Fig. 12. Slow strain rate testing for the studied alloy at a constant strain rate of 6.66 × 10-6 s-1: (a) (c) in air and (b) (d) in 0.35 wt.% NaCl solution. 22

ACCEPTED MANUSCRIPT Fig. 13. The relationship between loss factors of yield strength, ultimate tensile strength and tensile elongation under different aging treatment. Fig. 14. SEM fracture images of the SSRT samples aged for 24 h at different aging temperature failed in air: (a) UAI; (c) PA; (e) OAI; (g) OAII and failed in 3.5 wt.% NaCl solution: (b)

RI PT

UAI; (d) PA; (f) OAI; (h) OAII. Fig. 15. SEM fracture images of the SSRT samples aged at 120 °C for different durations failed in air: (a) UAII; (c)UAIII; (e) OAIII; (g) OAIV and failed in 3.5 wt.% NaCl solution: (b) UAII; (d) UAIII; (f) OAIII; (h) OAIV.

SC

Fig. 16. The relationship between SCC susceptibility and GBPs area fraction of the alloy under different treatments.

M AN U

Table Captions:

Table 1 Chemical compositions (wt.%) of the investigated alloy (Fe and Si, associated with the primary aluminum, are present as impurities) examined in the present investigation. Table 2 Heat treatment procedures of the investigated alloy.

Table 3 Chemical composition of the secondary phases in Fig. 1 (in at.%).

different temper.

TE D

Table 4 Summary of data showing grain boundary microstructures of the investigated alloy under

Table 5 SSRT results of the investigated alloy under different tempered conditions at a strain rate

AC C

EP

of 6.66 × 10-6 s-1.

23

ACCEPTED MANUSCRIPT Table 1 Chemical compositions (wt.%) of the investigated alloy (Fe and Si, associated with the primary aluminum, are present as impurities) examined in the present investigation. Zn

Mg

Mn

Cu

Sc

Zr

Si

Fe

Bal.

5.42

1.98

0.34

0.30

0.10

0.10

0.09

0.19

RI PT

Al

Table 2 Heat treatment procedures of the investigated alloy.

Solution (water quench) and aging treatment

USI

460 °C for 60 min

SC

Temper USII

470 °C for 20 min

470 °C for 40 min

M AN U

USIII PS

470 °C for 60 min

OSI

480 °C for 60 min

OSII

490 °C for 60 min

OSIII

470 °C for 90 min

OSIV

460 °C for 120 min

UAI

PS + 100 °C for 24 h PS + 120 °C for 4 h

TE D

UAII UAIII

PS + 120 °C for 12 h

PA

PS + 120 °C for 24 h PS + 140 °C for 24 h

OAII

PS + 160 °C for 24 h

OAIII

PS + 120 °C for 32 h

OAIV

PS + 120 °C for 48 h

AC C

EP

OAI

Table 3 Chemical composition of the secondary phases in Fig. 1 (in at.%).

Point

Al

Zn

Mg

Cu

Si

Fe

A

56.45

17.25

23.43

2.87

-

-

B

78.52

1.98

1.96

0.50

4.32

12.72

24

ACCEPTED MANUSCRIPT Table 4 Summary of data showing grain boundary microstructures1 of the investigated alloy under different temper. MPTs size (nm)

PFZ width (nm)

UAI

14.7 ± 3.0

11.3 ± 1.9

2.5 ± 0.4

20.7 ±1.8

PA

20.5 ± 3.2

19.8 ± 3.8

4.0 ± 0.5

29.5 ± 2.8

OAI

21.8 ± 3.4

21.0 ± 2.6

6.5 ± 1.0

40.9 ± 3.2

OAII

39.2 ± 6.8

23.3 ± 3.4

11.2 ± 2.6

49.8 ± 5.5

UAII

14.9 ± 2.8

11.6 ± 2.0

2.6 ± 0.4

UAIII

17.8 ± 2.8

14.0 ± 2.2

3.0 ± 0.6

OAIII

22.0 ± 3.2

21.8 ± 2.8

6.6 ± 1.0

OAIV 40.0 ± 7.0 23.8 ± 3.5 Data is the average values by statistics.

11.5 ± 2.8

RI PT

GBPs Af2 (%)

19.7 ± 1.5

21.5 ± 2.0 40.8 ± 3.4

48.1 ± 3.5

Af is the area fraction of GBPs.

M AN U

2

GBPs size (nm)

SC

1

Temper

Table 5 SSRT results of the investigated alloy under different tempered conditions at a strain rate of 6.66 × 10-6 s-1.

PA OAI

Air

470

3.5% NaCl

442

Air

518

3.5% NaCl

493

Air

466

3.5% NaCl

457

Air

461

AC C

OAII

Rm (MPa)

TE D

UAI

Corrosive environment

EP

Temper

UAII

UAIII

OAIII OAIV

3.5% NaCl

455

Air

450

3.5% NaCl

415

Air

472

3.5% NaCl

445

Air

468

3.5% NaCl

459

Air

460

3.5% NaCl

455

Rscc 6.0 4.8 1.9 1.3 7.8 5.7 2.0 1.1

25

Rp (MPa) 435 405 499 468 435 428 432 427 418 386 438 408 436 429 430 426

Rscc 6.9 6.2 1.6 1.1 7.6 6.8 1.6 0.9

A (%) 14.4 11.0 16.0 12.6 13.7 12.9 13.6 13.2 15.0 11.2 14.2 10.8 13.8 12.8 13.8 13.4

Rscc 23.6 21.2 5.8 5.7 25.3 24.0 7.3 2.9

ACCEPTED MANUSCRIPT

M AN U

SC

RI PT

(a)

B

EP

TE D

(b)

100 µm

AC C

A

Fig. 1. Microstructures of as-cast alloy: (a) OM image; (b) SEM image.

ACCEPTED MANUSCRIPT 0.4

Exo

0.0

RI PT

-0.2

Exo -0.4 -0.6

476.7 ºC

-0.8 -1.0 100

300

400

M AN U

200

SC

Heat flow/mV

0.2

Temperature/

AC C

EP

TE D

Fig. 2. DSC analysis of the Al-Zn-Mg-0.10%Sc-0.10%Zr alloy ingot.

500

(c)

(d)

SC

(b)

(f)

AC C

EP

(e)

TE D

M AN U

(a)

RI PT

ACCEPTED MANUSCRIPT

Fig. 3. Distribution of intergranular phases of different homogenization temperature for 24 h: (a) As-cast; (b) 460 °C; (c) 465 °C; (d) 470 °C; (e) 475 °C; (f) 480 °C.

(c)

(d)

SC

(b)

TE D

M AN U

(a)

RI PT

ACCEPTED MANUSCRIPT

100 µm

AC C

EP

Fig. 4. Distribution of intergranular phases for different homogenization time at 470 °C: (a) 8 h; (b) 16 h; (c) 24 h; (d) 32 h.

ACCEPTED MANUSCRIPT 2.0

(a) 480 1.5

1.0

465 0.5

RI PT

Heat flow/mV

470

460

As-cast 0.0

SC

-0.5

476.7

-1.0 200

300

400

2.0

(b) 1.5

1.0

24h

TE D

16h 0.5

8h

As-cast 0.0

-0.5

EP

Heat flow/mV

30h

500

M AN U

Temperature/

AC C

-1.0 200

476.7 300

400

500

Temperature/

Fig. 5. DSC curves of the alloy: (a) 24 h at different homogenization temperature; (b) 470 °C for different homogenization time.

(b)

(c)

(d)

(e)

(f)

(g)

(h)

RI PT

(a)

SC

ACCEPTED MANUSCRIPT

M AN U

(i)

50 µm

AC C

EP

TE D

Fig. 6. Microstructure of the studied alloys: (a) cold rolling; (b) USI; (c) PS; (d) OSI; (e) OSII; (f) USII; (g) USIII; (h) OSIII; (i) OSIV.

ACCEPTED MANUSCRIPT 600

Rm Rp A

(a)

12

500

450

400

RI PT

10

Elongation/%

Strength/MPa

550

8

460

470

480

490

Solution Temperature/

(b)

SC

600

Rm Rp A

450

400

20

M AN U

Strength/MPa

500

40

60

90

10

Elongation/%

12

550

8

120

TE D

Solution Time/min

AC C

EP

Fig. 7. Mechanical properties of the alloy under different solution treatment (aged at 120°C/24h): (a) soluted for 60 min at 460 °C, 470 °C, 480 °C, 490 °C; (b) soluted at 470 °C for 20 min, 40 min, 60 min, 90 min, 120 min. Rm, ultimate tensile strength; Rp, yield strength; A, elongation.

ACCEPTED MANUSCRIPT (a)

(b)

Rolling Direction

SC

RI PT

Al3(Sc,Zr)

(c)

M AN U

Al3(Sc,Zr)

(020)

TE D

(200)

AC C

EP

Fig. 8. Microstructures of the alloy after solution treatment at 470 °C for 60 min: (a) Fiber-like structure with subgrains, (b) Al3(Sc, Zr) within grains and in sub-grain boundaries; (c) SAD in [001]Al projection.

ACCEPTED MANUSCRIPT

600

Rm Rp A

(a)

14

550

10

450

6

100

120

140

Aging Temperature/ 600

Rm Rp A

550

500

TE D

Strength/MPa

160

M AN U

(b)

400

SC

8

400

450

Elongation/%

RI PT

500

4

12

16

14

12

10

Elongation/%

Strength/MPa

12

8

6 24

32

48

EP

Aging Time/h

AC C

Fig. 9. Mechanical properties of the alloy under different aging treatment after 470°C/60min solution treatment: (a) aged for 24 h at 100 °C, 120 °C, 140 °C, 160 °C; (b) aged at 120 °C for 4 h, 12 h, 24 h, 32 h, 48 h. Rm, ultimate tensile strength; Rp, yield strength; A, elongation.

ACCEPTED MANUSCRIPT (a)

(b)

(c)

Al3(Sc,Zr)

(e)

(f)

SC

(d)

RI PT

PFZ=40.4nm

(g)

(h)

M AN U

Al3(Sc,Zr)

(i)

TE D

Al3(Sc,Zr)

PFZ=48.7nm

AC C

EP

Fig. 10. Bright field microstructural evolution of the studied alloy under various aged conditions: (a) UAI; (b) PA; (c) OAI; (d) OAII; (e) UAII; (f) 120°C/24h, SAD in [011]Al projection; (g) UAIII; (h) OAIII; (i) OAIV.

ACCEPTED MANUSCRIPT 200

180 36

RI PT

35

140

34

120

Hardness Relative electric conductivity

SC

Hardness (HV)

160

Relative electric conductivity (% IACS)

37

80 0

10

M AN U

100

20

30

40

33

32 50

Aging time (h)

AC C

EP

TE D

Fig.11. Variation in the hardness and the relative electric conductivity of the alloy during aging process.

ACCEPTED MANUSCRIPT 600

UAI-Air

(a) 500

UAI-NaCl PA-NaCl

(b)

500

PA-Air OAI-Air

OAI-NaCl 400

OAII-Air

OAII-NaCl

Stress(MPa)

300

300

200

200

RI PT

Stress(MPa)

400

100

100

0

0 0

1

2

3

4

5

6

7

0

1

2

3

Displacement(mm)

5

6

(c)

(d)

500

500 400

300

200

M AN U

Stress(MPa)

400

UAIII-NaCl OAIII-NaCl OAIV-NaCl

300

200

100

100

0

0

1

2

3

4

5

6

7

8

TE D

Displacement(mm)

0

1

2

3

4

5

Displacement(mm)

EP

Fig. 12. Slow strain rate testing for the studied alloy at a constant strain rate of 6.66 × 10-6 s-1: (a) (c) in air and (b) (d) in 0.35 wt.% NaCl solution.

AC C

0

7

UAII-NaCl

SC

600

Stress(MPa)

4

Displacement(mm)

6

ACCEPTED MANUSCRIPT

(a)

28

Rm Rp A

6

24

20

4

16

RI PT

Rscc(Rm, Rp)/%

5

Rscc(A)/%

7

3

12

2

8

1

4

0 OAI

OAII

Aging Temper 9

(b)

Rm Rp A

7

Rscc(Rm, Rp)/%

6 5 4 3

0

TE D

2 1

28

M AN U

8

UAII

UAIII

24

20

16

Rscc(A)/%

PA

SC

UAI

12

8

4 OAIII

OAIV

Aging Temper

AC C

EP

Fig. 13. The relationship between loss factors of yield strength, ultimate tensile strength and tensile elongation under different aging treatment.

(b)

(c)

(d)

M AN U

SC

(a)

RI PT

ACCEPTED MANUSCRIPT

(f)

(h)

AC C

(g)

EP

TE D

(e)

Fig. 14. SEM fracture images of the SSRT samples aged for 24 h at different aging temperature failed in air: (a) UAI; (c) PA; (e) OAI; (g) OAII and failed in 3.5 wt.% NaCl solution: (b) UAI; (d) PA; (f) OAI; (h) OAII.

(b)

(c)

(d)

M AN U

SC

(a)

RI PT

ACCEPTED MANUSCRIPT

(f)

(h)

AC C

(g)

EP

TE D

(e)

Fig. 15. SEM fracture images of the SSRT samples aged at 120°C for different durations failed in air: (a) UAII; (c)UAIII; (e) OAIII; (g) OAIV and failed in 3.5 wt.% NaCl solution: (b) UAII; (d) UAIII; (f) OAIII; (h) OAIV.

ACCEPTED MANUSCRIPT 30

UAI PA

20

15

10

OAI

0 10

12

14

16

18

20

30

UAII

20

TE D

15

10

5

0

AC C

10

12

24

26

UAIII

EP

Loss factor of elongation,%

25

22

M AN U

GBPs Af

OAII

SC

5

RI PT

Loss factor of elongation,%

25

OAIII OAIV

14

16

18

20

22

24

GBPs Af

Fig. 16. The relationship between SCC susceptibility and GBPs area fraction of the alloy under different treatments.

ACCEPTED MANUSCRIPT Highlights: The suitable homogenization treatment of the alloy has been identified.




Evolution of microstructure and mechanical properties is investigated.




Strengthening mechanisms of the alloy has been established.




The basic mechanism has been proposed to model SCC of Al-Zn-Mg alloys.

AC C

EP

TE D

M AN U

SC

RI PT