materials Article

Precipitation Behavior and Quenching Sensitivity of a Spray Deposited Al-Zn-Mg-Cu-Zr Alloy Xiaofei Sheng 1,2 1 2 3

*

ID

, Qian Lei 1,3, *, Zhu Xiao 1 and Mingpu Wang 1

School of Materials Science and Engineering, Central South University, Changsha 410083, China; [email protected] (X.S.); [email protected] (Z.X.); [email protected] (M.W.) School of Materials Science and Engineering, Hubei University of Automotive Technology, Shiyan 442002, China Department of Materials Science and Engineering, College of Engineering, University of Michigan, Ann Arbor, MI 48109, USA Correspondence: [email protected]; Tel.: +1-734-763-5282

Received: 17 August 2017; Accepted: 17 September 2017; Published: 19 September 2017

Abstract: Precipitation behavior and the quenching sensitivity of a spray deposited Al-Zn-Mg-Cu-Zr alloy during isothermal heat treatment have been studied systematically. Results demonstrate that both the hardness and the ultimate tensile strength of the studied alloy decreased with the isothermal treatment time at certain temperatures. More notably, the hardness decreases rapidly after the isothermal heat treatment. During isothermal heat treatment processing, precipitates readily nucleated in the medium-temperature zone (250–400 ◦ C), while the precipitation nucleation was scarce in the low-temperature zone (400 ◦ C). Precipitates with sizes of less than ten nanometers would contribute a significant increase in yield strength, while the ones with a larger size than 300 nm would contribute little strengthening effect. Quenching sensitivity is high in the medium-temperature zone (250–400 ◦ C), and corresponding time-temperature-property (TTP) curves of the studied alloy have been established. Keywords: spray deposition; microstructure; aluminum alloys; strength; precipitation

1. Introduction In the past decades, 7xxx series (Al-Zn-Mg-Cu) aluminum alloys exhibited an outstanding performance and attracted much interest in many industry fields, especially in the automotive and the aviation field [1–3], because of their high strength, high ductility, and low density [4,5]. The high strength makes the products strong, and the high ductility allows the products to be easily forming during the machining process, and the low density saves the energy cost and reduces the carbon emission. Rolling, extrusion, drawing, and other conventional thermo-mechanical processing have limited capabilities to improve the strength of Al-Zn-Mg-Cu alloys [6–10]. To achieve a higher strength, some new technologies have been developed recently [11–14]. Severe plastic deformation (SPD) techniques, such as accumulative roll bonding [11], high-pressure torsion [12], equal channel angular pressing [13], and friction stir processing [14], were proven to be effective methods to refine the structure and enhance the strength. However, SPD techniques yield laboratory-scale samples, and they are still rare in industry [15]. Materials with fine grains usually exhibit a high strength (due to the Hall-Petch relation) [16–22]. The spray deposition technology has been developed rapidly [23,24]. Large-scale Al-Zn-Mg-Cu-Zr ingots with fine structure have been successfully fabricated by the spray deposition technology in our previous work [25,26]. The obtained ingots contain fine grains due to the rapid solidification and alloying [27–29]. Most investigations focus on the microstructure and properties of the spray-deposited Al-Zn-Mg-Cu alloy in the lab scale [26,30]. Thermal heat treatment is an efficient method to enhance the properties. However, some critical process parameters should Materials 2017, 10, 1100; doi:10.3390/ma10091100

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be reconsidered when the large ingots were turned to industrial products. For example, the solid solution-quenching treatment for precipitation hardening alloys is a very significant process, which would affect the followed aging process and the final properties [31]. In industry, ingots are about one ton in weight and a half cubic meter in volume, which are difficult to quench as fast as a sample with a lab-scale size. Decomposition will occur during the cooling down process. In addition, microstructure evolutions, precipitation behavior, properties variation, and quenching sensitivity of a spray deposited Al-Zn-Mg-Cu alloy during isothermal heat treatment (IHT) have not been reported. In this work, large-scale Al-Zn-Mg-Cu-Zr ingots were synthesized by a spray deposition technology. Microstructure evolution and properties variation of the solid solution treated alloy during isothermal heat treatments were investigated systematically. Precipitation behavior, quenching sensitivity, and strengthening mechanisms have been discussed. 2. Materials and Methods The studied alloy was synthesized with an SFZD-5000 type environmental chamber in Haoran company (Xinghua, China) [25,26]. The chemical compositions of the spray-deposited Al-Zn-Mg-Cu-Zr alloy are listed in Table 1. Before the spray deposition process, blocks of Al, Zn, Mg, Cu, and Al-Zr master alloy were melted, and then the molten melt was atomized by N2 gas. The distance of the atomizing deposition was 650 mm, rod ingots were spray deposition formed with a size of 500 mm in diameter and 1800 mm in length. The as-deposited ingot was extruded at 420 ◦ C with an extrusion speed of 3 mm/s and an extrusion ratio of 6.25:1 (500 mm to 200 mm in diameter). Experimental specimens were cut from the hot extruded rod, and they were solid solution treated at 470 ◦ C/1 h in the first furnace with a melted mixture of 50% NaNO3 and 50% KNO3 , and then rapidly transferred into another muffle furnace with a certain temperature (205–445 ◦ C) for isothermal heat treatment. Some of the isothermal heat treated specimens were aged at 120 ◦ C for 24 h. Table 2 gives the detailed parameters of the samples after different heat treatments. The sample A solid solution treated at 470 ◦ C/1 h was used as a reference set. Table 1. Chemical composition of studied Al-Zn-Mg-Cu-Zr alloy samples. Composition

Al

Zn

Mg

Cu

Zr

Fe

Si

Ti

Mn

Cr

wt % at %

Bal. Bal.

8.31 3.68

2.07 2.45

2.46 1.12

0.12 0.038

0.078 0.04

0.056 0.057

0.005 0.003

0.005 0.003

0.003 0.002

Table 2. Heat treatment parameters of the samples with different methods. Sample No. A B C D E F G H

Solution Treatment ◦ C/1

470 h 470 ◦ C/1 h 470 ◦ C/1 h 470 ◦ C/1 h 470 ◦ C/1 h 470 ◦ C/1 h 470 ◦ C/1 h 470 ◦ C/1 h

Isothermal Heat Treatment

Aging Treatment

/ / 205 ◦ C/5 s (30 min) 325 ◦ C/5 s (40 s), (30 min) 445 ◦ C/5 s, (30 min) 205 ◦ C/5 s (30 min) 325 ◦ C/5 s (40 s), (30 min) 445 ◦ C/5 s (30 min)

/ 120 ◦ C/24 h / / / 120 ◦ C/24 h 120 ◦ C/24 h 120 ◦ C/24 h

Hardness measurements of heat treated specimens were conducted on a HV-5 Vickers hardness tester (Laizhou Huayin Testing Instrument Co., Ltd., Laizhou, China) to keep track of the hardness variation with the aging time. The load is 2 kg, the dwell time is 15 s, and each average hardness was taken from at least seven measurements. Flat-shoulder tensile test samples under the same condition were prepared with dimensions of a gage length of 25 mm, and a width of 10 mm. Tensile testing experiments were performed on an MTS 810 material testing machine (MTS Systems Corporation, Eden Prairie, MN, USA) at room temperature with a strain rate of 2 mm/min. Microstructure

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observations on the isothermal heat treatment specimens were operated on an FEI Tecnai G2 F20 Materials 2017, 10, 1100 microscope (TEM) (FEI Company, Hillsboro, OR, USA) with an operation 3 of 14 transmission electron voltage of 200 kV. All TEM specimens were prepared by mechanical grinding, and then polished by a transmission electron microscope (TEM) (FEI Company, Hillsboro, OR, USA) with an operation double-jet electropolishing device (MTP-1A) (Shanghai Jiao Da Electrical and Mechanical Technology voltage of 200 kV. All TEM specimens were prepared by mechanical grinding, and then polished by ◦ C with a solution of 30% nitric acid +70% methanol. Development Co., Ltd., Shanghai, China) at −30 a double-jet electropolishing device (MTP-1A) (Shanghai Jiao Da Electrical and Mechanical Technology Development Co., Ltd., Shanghai, China) at −30 °C with a solution of 30% nitric acid +70% methanol.

3. Results

3. Results

3.1. Microstructure Evolutions during Isothermal Heat Treatment and Aging 3.1. Microstructure Evolutions during Isothermal Treatmentand andsegregation Aging For the as-spray deposited sample, there is Heat no dendrite due to rapid solidification, and smaller size were detected in thethere spray comparison For grain the as-spray deposited sample, is deposited no dendriteingot, and in segregation due to to that rapidof the solidification, and smaller grain size were detected in the spray deposited ingot, in comparison to conventional casting ingots [32]. Bright-field (BF) images and selected area diffraction pattern (SADPs) ◦ of thealloy conventional ingots [32].solid Bright-field (BF) images area quenched diffraction with of thethat studied (samplecasting A) having been solution treated at and 470 selected C/1 h and ◦ C/s of the studied (sample A) having been 1. solid solution at treatment, 470 °C/1 h and waterpattern with a(SADPs) quenching rate of 200alloy are shown in Figure After solid treated solution metallic quenched with water with a quenching rate of 200 °C/s are shown in Figure 1. After solid solution compounds precipitated during spray deposited process were re-dissolved into the matrix and a treatment, metallic compounds precipitated during spray deposited process were re-dissolved into supersaturated solid solution (SSS) formed. The SADP in Figure 1a shows that a set of diffraction spots the matrix and a supersaturated solid solution (SSS) formed. The SADP in Figure 1a shows that a set from of Aldiffraction were detected. The Al micrographs of the interiorofand boundary a SSS was spots from were detected. Thegrain micrographs the grain grain interior andshow grain that boundary formed after solid solution treatment (Figure 1b). show that a SSS was formed after solid solution treatment (Figure 1b).

(a)

(b)

Figure 1. Microstructure and selected area diffraction patter (SADP) of the sample having been

Figure 1. Microstructure and selected area diffraction patter (SADP) of the sample having been solution treated at 470 °C for 1 h. (a) SADP of the sample, beam along to zone axis of [100]Al; solution treated at 470 ◦ C for 1 h. (a) SADP of the sample, beam along to zone axis of [100]Al ; (b) bright-field image. (b) bright-field image.

The BF image and corresponding SADPs of the solid solution treated alloy (sample B) are shown in Figure 2. After aging treatment atSADPs 120 °C for 24 h, some precipitates were formed with B) a size The BF image and corresponding of the solid solution treated alloy (sample are of shown ◦ 10 nm (Figure 2b). Some diffraction spots were from the precipitates, as seen in Figure 2a. According in Figure 2. After aging treatment at 120 C for 24 h, some precipitates were formed with a size of 10 nm to the crystal geometry and crystal diffraction, indexing of the SADPs showed that the rod-like (Figure 2b). Some diffraction spots were from the precipitates, as seen in Figure 2a. According to the precipitates were η′-MgZn2 phase, which has hexagonal lattice parameters of a = 0.496 nm, c = 1.402 crystal geometry and crystal diffraction, indexing of the SADPs showed that the rod-like precipitates nm [33,34]. A series of crystal orientation relationship between the matrix and η′ phase could be werecalculated η0 -MgZn2and phase, which has hexagonal lattice parameters of a = 0.496 nm, c = 1.402 nm [33,34]. proved: (001) η′||(110)Al and [100]η′||[001]Al [35]. The high strength of Al-Zn-Mg0 phase could be calculated and A series of crystal orientation relationship between the matrix and η[36]. Cu alloys are due to a significant number of η′ phase precipitates Fine spherical L12-Al3Zr 0 0 proved: (001) η mixed ||(110)Al [100]η ||[001]Al [35]. The high1b, strength of Al-Zn-Mg-Cu alloys are precipitates with and rod-like η′-MgZn 2 precipitates in Figure and SADP in Figure 1a, proved due to a significant of η0after phase precipitates [36]. Fine spherical L12 -Al3 Zr precipitates mixed that the SSS wasnumber decomposed aging at 120 °C for 24 h. Figure 3 shows2 precipitates the BF images SADPs of the solid in solution alloy that having with rod-like η0 -MgZn in and Figure 1b, and SADP Figuretreated 1a, proved thebeen SSS was ◦ isothermal heat treated at 205 °C (quenching rate of 200 °C/s) for different durations (5 s and 30 min) decomposed after aging at 120 C for 24 h. (sample3 C). After isothermal heat treated 205solution °C for 5treated s, more diffraction spotsisothermal from Figure shows thebeing BF images and SADPs of the at solid alloy having been precipitates were in SADP (Figure 3a). The bright-field image of the grain inside showed G. heat treated at 205 ◦ Cdetected (quenching rate of 200 ◦ C/s) for different durations (5 s and 30 min) (sample C). P. zone and Al3Zr were formed (Figure 3b).◦ Increasing the IHT time to 30 min, the SADP in Figure 3c After being isothermal heat treated at 205 C for 5 s, more diffraction spots from precipitates were presents some diffraction spots from η′ precipitates, and the SSS was decomposed, and some detected in SADP (Figure 3a). The bright-field image of the grain inside showed G. P. zone and Al Zr precipitates with a size of 80 nm in length and 10 nm in diameter were detected (Figure 3d). Figure 4 3 wereshows formed 3b). the having IHT time 30 min, the SADP inatFigure BF(Figure images of theIncreasing studied alloys beentoisothermal heat treated 325 °C3cforpresents differentsome 0 precipitates, and the SSS was decomposed, and some precipitates with a diffraction spots from η durations (5 s, 40 s, and 30 min) (sample D). After isothermal heat treated at 325 °C for 5 s, η′

size of 80 nm in length and 10 nm in diameter were detected (Figure 3d). Figure 4 shows BF images

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of theMaterials studied alloys having been isothermal heat treated at 325 ◦ C for different durations4(5 s, 40 s, 2017, 10, 1100 of 14 ◦ C for 5 s, η0 precipitates were detected and 30 min) (sample D). After isothermal heat treated at 325 Materials 2017, 10, 1100 4 of 14 with precipitates a size of 150 nm detected in lengthwith anda20 nmofin150 diameter (Figure 4a). the IHT time4a). to 40 s, were size nm in length and 20 Increasing nm in diameter (Figure Increasing were the IHT time to 40 η′a precipitates were coarsened toand be sizenm of diameter 250 nm in(Figure length and precipitates were detected with nm in length and 20a100 nm in 4a). 4b). η0 precipitates coarsened tos, be asize sizeofof150 250 nm in length in diameter (Figure 100increasing nm in diameter (Figure 4b). thecoarsened IHTinside time tograin 30 η′ of precipitates grain Increasing the IHT time to 40 s,Further η′ precipitates were bemin, a coarsened size 250 nm length and Further the IHT time to 30 min,increasing η0 precipitates withinainside size of 400 nm coarsened with a size of 400 nm in length and 80 nm in diameter (Figure 4c). BF images and SADPs 100 nm in diameter (Figure 4b). Further increasing the IHT time to 30 min, η′ precipitates inside grain in length and 80 nm in diameter (Figure 4c). BF images and SADPs of the studied alloy having been of the studied alloy having been treated at diameter 445 °C for(Figure different durations (5 s and 30 min) coarsened with a size of 400 in length heat and 80 nm in 4c). BF images SADPs ◦ Cisothermal isothermal heat treated at 445 nm for different durations (5 s and 30 min) (sample E) are shown in (sample E) arealloy shown in Figure 5. After isothermal heat treatment at 445 °C for 5 s even 30 min, there of the studied having been isothermal heat treated at 445 °C for different durations (5 s and 30 min) Figure 5. After isothermal heat treatment at 445 ◦ C for 5 s even 30 min, there is few diffraction spots is few diffraction spots precipitates were detected in the SADP (Figure and there arethere few (sample E) are shown infrom Figure 5. After isothermal heat treatment at 445 °C for5a), 5 s even 30 min, from is precipitates were detected in the SADP (Figure 5a), and there are few precipitates (Figure 5b,c). precipitates (Figure 5b,c). few diffraction spots from precipitates were detected in the SADP (Figure 5a), and there are few precipitates (Figure 5b,c).

(a) (b) (a) (b) Figure 2. Bright-field (BF) images of the studied alloy having been solid solution treated at 470 °C for Figure 2. Bright-field (BF) images of the studied alloy having been solid solution treated at 470 ◦ C 1 h and then directly aged at 120 °C studied for 24 h. (a) having SADP,been the solid beamsolution along zone axis Al; Figure 2. Bright-field (BF) images of the alloy treated at of 470[100] °C for for 1 h and then directly aged at 120 ◦ C for 24 h. (a) SADP, the beam along zone axis of [100]Al ; image. aged at 120 °C for 24 h. (a) SADP, the beam along zone axis of [100]Al; 1(b)hbright-field and then directly (b) bright-field image. (b) bright-field image.

(a) (a)

(b) (b)

(c) (c)

(d) (d)

Figure 3. SADPs and BF images of the specimens having been isothermal heat treated at 205 °C for ; (b) treated the IHTattime is 5for s; different s and 30 min). SADP ofhaving (b), beam [110]Alheat Figure 3. durations SADPs and(5BF images of the(a) specimens beenalong isothermal 205 °C Figure 3. SADPs and BF images of the specimens having been isothermal heat treated at 205 ◦ C for (c) SADP durations of (d), beam Al; and(d) the isothermal heatalong treatment is 30 time min. is 5 s; Al; (b) time the IHT different (5 along s and [110] 30 min). (a) SADP of (b), beam [110](IHT)

different durations (5 s and 30 min). (a) SADP of (b), beam along [110]Al ; (b) the IHT time is 5 s; (c) SADP of (d), beam along [110]Al; and(d) the isothermal heat treatment (IHT) time is 30 min. (c) SADP of (d), beam along [110]Al ; and(d) the isothermal heat treatment (IHT) time is 30 min.

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(a) (a)

(b) (b)

(c) (c) Figure 4. BF images of the studied alloy having been isothermal heat treated at 325 °C for different Figure 4. BF images of the studied alloy having been isothermal heat treated at 325 ◦ C for different Figure 4. BF images of the studied alloy having durations. The IHT durations are: (a) 5 s; (b) 40 s;been and, isothermal (c) 30 min. heat treated at 325 °C for different durations. The IHT durations are: (a) 5 s; (b) 40 s; and, (c) 30 min. durations. The IHT durations are: (a) 5 s; (b) 40 s; and, (c) 30 min.

(a) (a)

(b) (b)

(c) (c)

(d) (d)

Figure 5. SADPs and BF images of the studied alloy having been isothermal heat treated at 445 °C for 5. and BF images of the alloy been isothermal heattime treated °C for 5Figure 30SADPs min. (a) SADP from the areastudied in (b); (b) thehaving isothermal treatment is 5ats;445 (c) SADP Figures5.and SADPs and BF images of the studied alloy having beenheat isothermal heat treated at 445 ◦ C for 5 s and 30 min. (a) SADP from the area in (b); (b) the isothermal heat treatment time is 5 s; (c) SADP from the area in (d); and (d) BF image, the isothermal heat treatment time is 30 min. 5 s and 30 min. (a) SADP from the area in (b); (b) the isothermal heat treatment time is 5 s; (c) SADP from the area in (d); and (d) BF image, the isothermal heat treatment time is 30 min.

from the area in (d); and (d) BF image, the isothermal heat treatment time is 30 min.

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Figure 66 shows shows BF BF images images and and SADPs SADPs of of the the solution solution treated treated alloy alloy having having been been isothermally isothermally heat heat Figure ◦ C for different durations (5 s and 30 min), and then aged at 120 ◦ C for 24 h (sample F). treated at 205 treated at 205 °C for different durations (5 s and 30 min), and then aged at 120 °C for 24 h (sample F). ◦ C for 5 s and aged at 120 ◦ C for 24 h, some diffraction spots from After After the the sample sample was was treated treated at at 205 205 °C for 5 s and aged at 120 °C for 24 h, some diffraction spots from 0 η and Al Zr precipitates were detected in the the SADP SADP (Figure (Figure 6a), 6a), and and the the bright-field bright-field images images of of grain grain η′ and Al33Zr precipitates were detected in 0 precipitates, whose size was 20–40 nm (Figure 6b). inside and grain boundary also shows many η inside and grain boundary also shows many η′ precipitates, whose size was 20–40 nm (Figure 6b). Increasing the the IHT IHT time time to to 30 Increasing 30 min, min, the the SADP SADP in in Figure Figure 6c 6c presented presented many many diffraction diffraction spots spots from from 0 precipitates with a size of 100 nm in length and 10 nm in diameter precipitates, and many coarse η precipitates, and many coarse η′ precipitates with a size of 100 nm in length and 10 nm in diameter were were detected detected (Figure (Figure 6d). 6d).

(a)

(b)

(c)

(d)

Figure heat treated treated at at 205 205 ◦°C Figure 6. 6. SADPs SADPs and and BF BF images images of of the the studied studied alloy alloy having having been been isothermal isothermal heat C for for different durations (5 s and 30 min), and then aged at 120 °C for 24 h. (a) SADP from the area in (b); ◦ different durations (5 s and 30 min), and then aged at 120 C for 24 h. (a) SADP from the area in (b); (b) the theisothermal isothermalheat heat treatment (c) SADP from theinarea in (d); and (d) isothermal heat (b) treatment is 5iss; 5(c)s;SADP from the area (d); and (d) isothermal heat treatment treatment is 30 min. is 30 min.

Bright-field TEM images and SADPs of the solution treated alloy having been annealed at 325 ◦°C Bright-field TEM images and SADPs of the solution treated alloy having been annealed at 325 C for different durations (5 s, 40 s, and 30 min), and then aged at 120 °C for 24 h (Sample G) are shown for different durations (5 s, 40 s, and 30 min), and then aged at 120 ◦ C for 24 h (Sample G) are shown in in Figure 7. After isothermal heat treated at 325 °C for 5 s and aging at 120 °C for 24 h, η′ precipitates Figure 7. After isothermal heat treated at 325 ◦ C for 5 s and aging at 120 ◦ C for 24 h, η0 precipitates were were formed with a size of 10–20 nm (Figure 7a). Increasing the IHT time to 40 s, coarse η′ precipitates formed with a size of 10–20 nm (Figure 7a). Increasing the IHT time to 40 s, coarse η0 precipitates were were formed with a length of 300 nm and a diameter of 30 nm (Figure 7b). Increasing the IHT time to formed with a length of 300 nm and a diameter of 30 nm (Figure 7b). Increasing the IHT time to 30 min, 300 min, η′ precipitates with a size of 300 nm in length and 100 nm in diameter (Figure 7c). Figure 8 η precipitates with a size of 300 nm in length and 100 nm in diameter (Figure 7c). Figure 8 shows shows BF images of the solution treated alloy having been isothermal heat treated at 445 °C for BF images of the solution treated alloy having been isothermal heat treated at 445 ◦ C for different different durations (5 s and 30 min), and then aged at 120 °C for 24 h (sample H). After isothermal durations (5 s and 30 min), and then aged at 120 ◦ C for 24 h (sample H). After isothermal heat treated heat treated at 445 °C for 5 s and aged at 120 °C for 24 h, some η′ precipitates were detected with a at 445 ◦ C for 5 s and aged at 120 ◦ C for 24 h, some η0 precipitates were detected with a size of 10–20 nm size of 10–20 nm (Figure 8a), and they maintained same size even if isothermal heat treated for 30 min (Figure 8a), and they maintained same size even if isothermal heat treated for 30 min (Figure 8b). (Figure 8b). Al3Zr precipitates were also detected in these specimens (Figures 7 and 8). Al3 Zr precipitates were also detected in these specimens (Figures 7 and 8).

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(a) (a)

(b) (b)

(c) (c) Figure 7. BF images of the studied alloy having been isothermal heat treated at 325 °C for different Figure 7. BF images of the studied alloy having been isothermal heat treated at 325 ◦ C for different durations (5 images s, 40 s, and 30studied min), and then agedbeen at 120 °C for 24heat h. The isothermal heat Figure 7. BF of the alloy having isothermal treated at 325 °C fortreatment different durations (5 s, 40 s, and 30 min), and then aged at 120 ◦ C for 24 h. The isothermal heat treatment durations are: (b) 40 and, and (c) 30then min.aged at 120 °C for 24 h. The isothermal heat treatment (5 s,(a) 405s,s;and 30s;min), durations are: (a) 5 s; (b) 40 s; and, (c) 30 min. durations are: (a) 5 s; (b) 40 s; and, (c) 30 min.

(a) (a)

(b) (b)

Figure 8. BF images of the studied alloy having been isothermal heat treated at 445 °C for different durations (5 simages and 30of min), thenalloy agedhaving at 120 °C forisothermal 24 h. The isothermal heat Figure 8. BF the and studied been heat treated at treatment 445 °C fordurations different Figure 8. BF images of the studied alloy having been isothermal heat treated at 445 ◦ C for different are: (a) 5 s;(5(b) 30 min. durations s and 30 min), and then aged at 120 ◦°C for 24 h. The isothermal heat treatment durations durations (5 s and 30 min), and then aged at 120 C for 24 h. The isothermal heat treatment durations are: (a) 5 s; (b) 30 min.

are: (a) 5 s; (b)Variation 30 min. 3.2. Hardness 3.2. Hardness Variation The effect of IHT on the mechanical properties of the alloy was investigated through hardness 3.2. Hardness Variation The effect of IHT on the mechanical properties of the alloy investigated through measurements. Figure 9 shows the hardness variation versus thewas isothermally heat treatedhardness time on measurements. Figure 9 shows the solution hardness variation versus the 1was isothermally heat treated time on studied alloy been treated at for h and quenched (cooling rate of The effect of samples IHT on having the mechanical properties of470 the°C alloy investigated through hardness studied alloy samples having beenhardness at 470 °C 1 hisothermally and quenchedand (cooling ratetime of 200 °C/s), isothermal treated atsolution differenttreated temperatures forfor various durations, then aged at on measurements. Figure 9heat shows the variation versus the heat treated 200 °C/s), isothermal heat treated at different temperatures for various durations, and then aged 120 °C for 24 h. With increasing the IHT time, the hardness decreased successively in all IHT ◦ studied alloy samples having been solution treated at 470 C for 1 h and quenched (cooling at rate of 120 °C for 24 Especially h. With increasing IHT time, the hardness decreased in all temperatures. when the the samples were isothermal heat treated atsuccessively 295 °C (or 325 °C)IHT for 200 ◦ C/s), isothermal heat treated at different temperatures for various durations, and then aged Especially when therapidly samples were isothermal at 295 °C) for 3temperatures. min, the hardness values were decreased from 210heat HV treated to 110 HV for °C 295(or °C,325 and from at 120 ◦ C for 24 h. With increasing the IHT time, the hardness decreased successively in all IHT 3 min, hardness values were rapidly decreased fromisothermal 210 HV toheat 110treated HV for °C,(sample and from 210 HV the to 120 HV for 325 °C. However, when samples were at 295 205 °C F), ◦ C (or 325 ◦ C) for temperatures. Especially theThe samples were isothermal heatHV treated 295 210 HV to 120 HV for 325when °C. However, when samples were isothermal heat at 205 °C (sample F), the hardness decreased slightly. hardness decreased from 210 totreated 198 at HV when the sample ◦ C, and from 210 HV 3 min,was the hardness values were decreased from a210 HV 210 to 110 for HV 295 the hardness decreased slightly. The from HVHV to 198 when isothermal heat treated atrapidly 205 °C hardness for 3 min.decreased When specimen was solution treated atthe 470sample °C for to 120 HV for 325 ◦heat C. However, when samples were aisothermal heat treated at 205at◦470 C (sample was isothermal treated at 205 °C for 3 min. When specimen was solution treated °C for F),

the hardness decreased slightly. The hardness decreased from 210 HV to 198 HV when the sample was isothermal heat treated at 205 ◦ C for 3 min. When a specimen was solution treated at 470 ◦ C for

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◦C 111 hh hardness was was as as high high as as 210 210 HV. HV. and then immediately h and and then then immediately immediately aged aged at at 120 120 °C °C for for 24 24 h h (sample (sample A), A), the the hardness hardness was as high as 210 HV. The hardness of specimens decreased slightly from 210 HV to 202 HV when the SSS specimens were The hardness of specimens decreased slightly from 210 HV to 202 HV when the SSS specimens were The hardness of specimens decreased slightly from 210 HV to 202 HV when the SSS specimens were ◦ C and kept for 3 min, followed by aging at 120 ◦ C transferred at at thethe temperature of 445 transferred into furnace temperature of °C transferredinto intoa aafurnace furnace at the temperature of 445 445 °C and and kept kept for for 33 min, min, followed followed by by aging aging at at for 24 h (sample H). As samples were transferred into the IHT furnace, kept warm in the furnace 120 °C for 24 h (sample H). As samples were transferred into the IHT furnace, kept warm the 120 °C for 24 h (sample H). As samples were transferred into the IHT furnace, kept warm in in for the ◦ C for 24 h, the hardness of specimens was less than 160 HV. 30 min, and then followed by followed aged at 120 furnace for min, and by furnace for 30 30 min, and then then followed by aged aged at at 120 120 °C °C for for 24 24 h, h, the the hardness hardness of of specimens specimens was was less less Further increasing the IHT time, hardness decreased successively. than 160 HV. Further increasing the IHT time, hardness decreased successively. than 160 HV. Further increasing the IHT time, hardness decreased successively.

Figure 9. Hardness variation variation via via the the isothermal isothermal treated time time on the studied samples having been Figure Figure 9. 9. Hardness isothermal treated time on on the the studied studied samples samples having having been been ◦ solution treated at 470 °C for 1 h, isothermal heat treated for various durations, and then aged solution forfor 1 h,1 isothermal heatheat treated for various durations, and then 120 ◦ at C solutiontreated treatedatat470 470C°C h, isothermal treated for various durations, andaged thenataged at 120 °C for 24 h. for 24 h. 120 °C for 24 h.

3.3. Stress-Strain Curves 3.3. 3.3. Stress-Strain Stress-Strain Curves Curves Figure 10 shows the tensile stress-strain curves of the tensile samples. For the solid solution Figure Figure 10 10 shows shows the the tensile tensile stress-strain stress-strain curves curves of of the the tensile tensile samples. samples. For For the the solid solid solution solution treated sample directly aged at 120 °C for 24 h (sample A), the ultimate tensile strength was 719.8 ◦ treated C for 24 h (sample A), the ultimate tensile strength was 719.8 MPa. treated sample sample directly directly aged aged at at 120 120 °C MPa. While for samples having been isothermal heat treated at 205 °C, 325 °C, 445 °C, then aged at 120 ◦ ◦ ◦ While C, 325 C, 445 C, then C While for for samples samples having havingbeen beenisothermal isothermalheat heattreated treatedat at205 205 °C, 325 °C, 445 °C, then aged aged at at 120 120 ◦°C °C for 24 h, the ultimate tensile strengths were 620.2 MPa (for 205 °C, sample F/5 s), 606.4 MPa (for 325 °C, ◦ for h, the theultimate ultimatetensile tensile strengths were 620.2 MPa C, sample s),MPa 606.4(for MPa for 24 h, strengths were 620.2 MPa (for(for 205 205 °C, sample F/5 s),F/5 606.4 325(for °C, sample MPa (for sample H/5 s), While been ◦ C, G/5 325 sample G/5647.5 s), and MPa°C, (for 445 ◦ C, sample H/5 s), respectively. While having for samples sample G/5 s), s), and and 647.5 MPa647.5 (for 445 445 °C, sample H/5 s), respectively. respectively. While for for samples samples having been isothermal heat treated at 205 °C, 325 °C, 445 °C for 30 min, and then aged at 120 °C for 24 h, ◦ ◦ ◦ ◦ having beenheat isothermal treated at 205 C, 325 C, 30 445min, C for 30then min,aged and then aged C the for isothermal treated heat at 205 °C, 325 °C, 445 °C for and at 120 °C at for120 24 h, the ultimate tensile strength values were 524.5 MPa (for 205 °C, sample G/30 min), 300.1 MPa (for 325 °C, ◦ 24 h, the ultimate tensile values strength values were MPa°C, (for 205 C, sample 300.1 MPa ultimate tensile strength were 524.5 MPa524.5 (for 205 sample G/30 min),G/30 300.1 min), MPa (for 325 °C, sample G/30 min), and 647.5 MPa (for 445 °C, sample G/30 min). Table 3 shows the detailed sample min), G/30 and 647.5 (for 445 min). Table 3 shows detailed (for 325 ◦G/30 C, sample min), MPa and 647.5 MPa°C, (forsample 445 ◦ C,G/30 sample G/30 min). Table 3the shows the mechanical properties of with heat (see 2). mechanical properties of the the specimens specimens with different different heat treatments treatments (see Table Table 2).Table 2). detailed mechanical properties of the specimens with different heat treatments (see

(a) (a)

(b) (b)

Figure 10. Stress-strain curves of the tensile samples having been isothermal heat treated at 205 Figure 10. 10. Stress-strain Stress-strain curves curves of of the the tensile tensile samples samples having having been been isothermal isothermal heat heat treated treated at at 205 205 ◦°C, °C, Figure C, 325 °C, and 445 °C for different durations (5 s and 30 min), then aging at 120 °C for 24 h. (a) the 325 °C, and 445 °C for different durations (5 s and 30 min), then aging at 120 °C for 24 h. (a) the ◦ ◦ ◦ 325 C, and 445 C for different durations (5 s and 30 min), then aging at 120 C for 24 h. (a) the isothermal heat treated time is (b) the isothermal heat treated time 30 min. isothermal heat heat treated treated time time is is 555s;s; s;(b) (b)the theisothermal isothermalheat heattreated treatedtime timeisis is30 30min. min. isothermal

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Table 3. Mechanical properties of the studied specimens subjected to isothermal heat treatments for Table 3. Mechanical properties of the studied specimens subjected to isothermal heat treatments for different treatment parameters. different treatment parameters.

Sample No. Sample B No. C/5 B s C/5 ss D/5 D/5 s E/5 s E/5 s F/30 F/30 min min G/30 G/30 min min H/30 min H/30

Elongation (%) Ultimate Tensile Strength (MPa) Yield Strength (MPa) Elongation (%) Ultimate Tensile Yield Strength 7.7 719.8 Strength (MPa) 699.8 (MPa) 6.97.7 620.2 598.2 719.8 699.8 620.2 598.2 7.36.9 604.4 578.3 604.4 578.3 7.77.3 647.5 621.9 7.7 647.5 621.9 7.77.7 524.5 506.4 524.5 506.4 8.98.9 300.1 267.6 300.1 267.6 647.5 621.2 7.37.3 647.5 621.2

4. Discussion 4.1. The Quenching Sensitivity Temperature-time-property (TTP) Temperature-time-property (TTP) curves curves can can be be established established by the hardness of the samples with isothermal heat treatment and aging treatment. The The nucleation nucleation rate rate (k) (k) can can be expressed with the isothermal following equation [37,38]:

k3kk34k24 2 k5 k tctc =k exp[ ]exp[ k =k = = k22 exp[ ] exp[ ]5 ] 2 2 k RT k  T RT ( ) k1 1 RT RT (k4 4 − T )

(1) (1)

where t tcisisthethe critical required to precipitate a constant amount is the IHT where critical timetime required to precipitate a constant amount of solutes;ofT solutes; is the IHTTtemperature c − 1 −1 temperature K).molar R is the gas molar k1, kk52,are k3, materials k4, and k5constants are materials (degree K). R(degree is the gas constant (8.31 Jconstant ·K ). k1 ,(8.31 k2 , k3J·K , k4 , ).and that constants that are related to the natural logarithm of the volume fraction of untransformed are related to the natural logarithm of the volume fraction of untransformed precipitates (k1 ), the precipitates 1), number the reciprocal of the sites number of nucleation sites (kof 2), nucleation the formation of reciprocal of (k the of nucleation (k2 ), the formation energy sites energy (k3 ), solid nucleation sites (k3), solid temperature (k4), and the diffusion energy (k Thecurves fitted solution temperature (k4 ), solution and the diffusion activation energy (k5 ). Theactivation fitted coefficients of5).TTP −11 for k data coefficients of TTP curves from the Figure by the areJ/mol 2.10 ×for 10k−11 2, 1061 from the Figure 11 by the hardness data are112.10 × 10hardness 802kK for k4J/mol , and 2 , 1061 3 , for 5 5 for k× 3, 802 K for kfor 4, and 1.23 × 10 J/mol k5, respectively. TTP curves aretemperatures “C” shape, whose nose 1.23 10 J/mol k5 , respectively. TTPfor curves are “C” shape, whose nose are around temperatures 325 ◦ C (Figureare 11).around 325 °C (Figure 11).

Figure 11. The time-temperature-property (TTP) been Figure 11. The time-temperature-property (TTP) curves curves of of the the studied studied alloy alloy subjected subjected to to have have been ◦ solid solution treated at 470 °C for 1 h, isothermal heat treatment for various durations, and then aged solid solution treated at 470 C for 1 h, isothermal heat treatment for various durations, and then aged ◦ C for at at 120 120 °C for 24 24 h. h.

For the studied alloy, we divided the temperature zone during the quenching process into lowFor the studied alloy, we divided the temperature zone during the quenching process into temperature zone (400 C). When the supersaturated solid solution were isothermal heat treated at the low-temperature temperature◦ zone (400 °C). The hardness decreased slightly. zone (400 C). The hardness decreased slightly. While, if SSS While, if SSS samples were isothermal heat treated at the medium-temperature zone (250–400 °C), the hardness decreased rapidly, indicating that the alloy presented a high quenching sensitivity and

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samples were isothermal heat treated at the medium-temperature zone (250–400 ◦ C), the hardness decreased rapidly, indicating that the alloy presented a high quenching sensitivity and a short incubation period in the medium-temperature zone. Similar TTP curves of conventional casting 7XXX alloy have been reported by some researchers, and many methods were applied to study the quenching sensitivity. In-situ electrical resistivity measurement and differential calorimetry were employed to investigate the high temperature precipitation kinetics and TTT curves, which reveal that the nose temperature decreased from 375 ◦ C to 350 ◦ C when the precipitated fraction increase [37]. Liu et al. have investigated the TTP diagram for Al-Zn-Mg-Cu alloys including 7075, 7175, 7050, 7010, 7055, 7085, and 1933. Their result demonstrated that the nose temperature of TTP diagrams changed with the chemical composition of alloys, which was 355 ◦ C for 7055 alloy while it was 295 ◦ C for 7085 alloy [39]. Nie et al. reported that the aging hardness reductions for the 7085, 7050, 7150, and 7055 alloys by air quenching are up to 8.6%, 18.4%, 21%, and 21.7%, respectively, in comparison that that by water quenching [40]. Air quenching presents lower a quenching rate when compared with water quenching. The aging yield strength reductions for the studied alloy are 27.6%, 31.4%, and 11.2% for sample F/5 s, sample G/5 s, and sample H/30 min, respectively, in comparison that with water quenching, thus a higher furnace-to-furnace transfer temperature is preferred for the studied alloy. Due to recrystallization, Zr-containing dispersoids become incoherent with the matrix and act as effective nucleation sites for heterogeneous precipitation during slow quenching, thus Zr addition in the alloy inhibits recrystallization that is sensitive to the quench sensitivity [41]. The hardness of 7085 alloy increased with an increase of Mg, which was reported by Deng et al. who have studied the quench sensitivity of Al alloys via depth of age hardening layer according to hardness retention value [27]. The TTP curve of 7010 alloy indicated that transformation occurred most rapidly between 250 and 400 ◦ C [42]. For the above 7XXX alloy, the nose temperature of TTP curves were between 250–400 ◦ C, which are consistent with the result in this work (325 ◦ C). As the SSS was isothermal heat treated, some precipitates have been detected in above images and previous investigations [37,43,44]. High quenching sensitivity is the result of fast precipitation from the supersaturated solid solution [24]: • • 4 π • • dVt = π ·V0 · N · G3 ·t3 · exp(− G3 Nt4 ) dt 3 3 •

(2) •

where V t is the precipitates volume, V 0 is the sample volume, N is the nucleating rate of precipitates, G3 is the precipitate growing rate, and t is IHT time. When the SSS samples were isothermal heat treated •

•

in the medium temperature zone (250–400 ◦ C), N · G3 reached the maximum value. The incubation period of the studied alloy is short in medium temperature zone but long in the low-temperature zone and high-temperature zone. Thus, for the studied alloy, after solid solution treatment, the samples should be rapidly transferred into the aging treatment furnace, and the temperature of the sample during transfer process should be higher than 400 ◦ C. Meanwhile, if the sample could not be moved to the aging furnace very quickly, a water quenching should be performed for cooling down the sample temperature to lower than 250 ◦ C. Otherwise, the SSS sample would be decomposed and coarse precipitates would form on grain boundary and grain inside, resulting in undesired mechanical properties. 4.2. Precipitation Behavior during Isothermal Heat Treatment The studied alloy was a precipitation hardening aluminum alloy. After the appropriate solution annealing and quenching, a supersaturated solid solution would be formed; then decomposition behavior would start from quenching, as well as aging at a low temperature. Previous investigations proved that the SSS would begin to decompose at 120 ◦ C [45]. Generally, for the Al-Zn-Mg-Cu-Zr alloy, it will take 12 h to finish the peak aging process [43]. The precipitates size was only 5–8 nm in length and 1.5–3 nm in width. With increasing the aging temperature, the decomposition behavior

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accelerated [46]. During the isothermal heat treatment processing, few η0 precipitates were detected when the IHT time is only 5 s (Figure 3b). Increasing the time to IHT 30 min, some precipitates with a size of 80 nm in length and 10 nm in diameter were detected (Figure 3d). However, when the SSS was quenched to isothermal temperature of 325 ◦ C (Figure 4), even the IHT time is 5 s, some η0 precipitates were detected with a size of 150 nm in length and 20 nm in diameter (Figure 4a). Similarly, the SSS sample was quenched to an isothermal temperature of 445 ◦ C, precipitation happened (Figure 5). These isothermal heat treated specimens were aging treated at 120 ◦ C for 24 h, many η0 -MgZn2 were precipitated from the matrix. As a result of this work, an appropriate isothermal temperature required to be tailored, low-temperature zone (400 ◦ C) were worked out for the studied alloy to maintain small precipitates and high strength. 4.3. Influence of Precipitates Size on the Strengthening Response Low-temperature aging treatment at 120 ◦ C contributed to the precipitation hardening in the studied alloy. However, the precipitation occurred in samples during transferring from the solution treated furnace to the aging furnace, and the precipitates may coarsen. The Orowan mechanism for dislocation and particles can evaluate the significant increase of yield strength by the precipitates. The increment of ∆σOrowan can be expressed as follows [36].
ln dp /b
∆σ = 2π(1 − υ)1/2 λ − dp 0.81MGb

1 λ = dp 2

s

3π 2 fv

(3)

(4)

where, M is the Taylor-factor (M for Al is 3.06); G is the shear modulus (shear modulus for Al is 26.9 GPa); dp is the average diameter of precipitates particle; b is the Burgers vector (b for Al is 0.286 nm); ν is the Poisson's ratio (ν for Al is 0.33) [47]; λ is the average particle plane square lattice spacing; f v is the volume fraction of particles. To simplify the calculation, we assumed that all of the Mg and Zn atoms precipitated in the form of η0 -MgZn2 precipitates [48]. The volume fraction of MgZn2 precipitates can be calculated and f v is 5.9%. Therefore, the increment of ∆σOrowan was a function of dp and can be calculated as follows: ∆σOrowan =

1286 + 1029 ln dp dp

(5)

When the sample was solution-treated and immediately aged at 120 ◦ C for 24 h, dp was about 10 nm (Figure 2, sample B), the ∆σOrowan was 365.5 MPa. However, when the dp was about 300 nm in Figure 7c (the alloy was isothermal treated at 325 ◦ C for 30 min then aged at 120 ◦ C for 24 h, sample G), the ∆σOrowan is only 23.9 MPa. The precipitate size is a significant factor on the ∆σOrowan strengthening effect. 5. Conclusions In summary, a high strength Al-Zn-Mg-Cu-Zr alloy was synthesized by using a spray deposition technology followed by hot extrusion, solution treatment, isothermal heat treatment, and aging treatment. (1)

(2)

The hardness and the ultimate tensile strength decreased with the isothermal heat treatment time. TTP curves of the studied alloy have been established, and the mechanism for high quenching sensitivity at the nose temperature zone has been discussed. TTP curves of the studied alloy are “C”-like ones, and the nose-tip temperatures of these curves are around 325 ◦ C. The studied alloy presented a high quenching sensitivity at the medium-temperature zone (250–400 ◦ C) because in this medium temperature zone the nucleation rate of the precipitates was

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high and a significant number of precipitates formed. After solid solution treatment, precipitation appeared scarcely as specimens had been isothermal heat treated at the low-temperature zone (400 ◦ C), both of above two temperature zones were available for the industry production. The strengthening phase would precipitate as the sample aged at 120 ◦ C, resulting in high mechanical properties. Precipitates with size of 10 nm would contribute a significant increase in yield strength, while ones larger than 300 nm contribute only little.

Acknowledgments: The authors thank to the financial support by the National Key Technology K & D Program (2014BAC03B08), the National Key Research and Development Program of China (2016YFB0301300), the start-up funding from the Central South University, the Fundamental Research Funds for the Central Universities of Central South University (2014ZZTS018) and Hunan Provincial Innovation Foundation for Postgraduate (CX2014B048). Author Contributions: Xiaofei Sheng and Qian Lei conceived and designed the experiments; Xiaofei Sheng performed the experiments; Xiaofei Sheng and Qian Lei analyzed the data; Mingpu Wang and Zhu Xiao contributed materials and discussions; Xiaofei Sheng and Qian Lei wrote the paper. Conflicts of Interest: The authors declare no conflict of interest.

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