key: cord-0303310-loimvolo
authors: Pan, Shuaihang; Yao, Gongcheng; Yuan, Jie; Sokoluk, Maximilian; Li, Xiaochun
title: Manufacturing of Bulk Al-12Zn-3.7Mg-1Cu Alloy with TiC Nanoparticles
date: 2020-12-31
journal: Procedia Manufacturing
DOI: 10.1016/j.promfg.2020.05.054
sha: 7dda2ad6eb4d0f3a55e01ff59fb3b3d4be0da149
doc_id: 303310
cord_uid: loimvolo

Abstract Aluminum alloys are significant in various applications, and the desire is strong to develop higher strength with balanced other properties out of high Zn and Mg contents in aluminum. However, high contents of Zn and Mg will make the matrix brittle and incompatible for mass manufacturing. As nanoparticles could effectively and dynamically control the brittle secondary phases, it opens a new route to design and manufacture high-Zn and high-Mg aluminum alloys. Here, we demonstrate the feasibility of nanomanufacturing of Al-12Zn-3.7Mg-1Cu alloy system with the help from TiC nanoparticles for potential ultrahigh-strength light-weight alloys. By applying the “hot rolling-natural aging” and “solution treatment-natural aging” processing routes, Al-12Zn-3.7Mg-1Cu alloy with TiC nanoparticles can achieve higher microhardness via natural aging and better machinability when compared with the pure Al-12Zn-3.7Mg-1Cu alloy. Microstructure and thermodynamic studies reveal that the reasons for the peak hardness difference of Al-12Zn-3.7Mg-1Cu alloy and Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC nanocomposites are due to the free energy difference by solution treatment temperature and the strain energy sustained by the residual secondary phases. The enhanced mechanical properties in Al-12Zn-3.7Mg-1Cu alloy with TiC nanoparticles at the peak-microhardness state after “solution treatment-natural aging” processing route prove the feasibility of the abovementioned nano-treating manufacturing method.

Al-Zn-Mg-Cu alloys (7xxx series aluminum alloys) are extensively studied and used for aeronautical applications where high strength and toughness are required [1] . These alloy systems are mainly composed of easily procured elements of Zn, Mg, and Cu, and the high strength with low cost for raw materials makes these alloys irreplaceable. It has been known that three ways are effective to optimize the properties of Al-Zn-Mg-Cu alloy: 1) Changing the main alloying elements of Zn, Mg, and Cu; 2) Tuning the trace elements (Sc, Cr, Zr, and Mn); 3) Minimizing the atomic-scale element impurity (Fe, Si) [2] . All these methods are important for achieving a suitable combination of precipitate formation like η(MgZn2), T(Al2Mg3Zn3), S(Al2CuMg) and θ(Al2Cu).

Recently, with strong needs for higher strength and balanced other properties, to mitigate the costs by the expensive trace elements, research on Al alloys with higher-Zn (>8.5 wt. %) and higher-Mg becomes a hot topic [3] - [6] . Certain theoretical calculations linking secondary phase show the promising future of using high-concentration Zn and Mg for Al alloy strengthening [7] . However, with more Zn and Mg into the aluminum alloy system, the alloy may be difficult to deform due to the brittle large precipitates. Without effective modifications for these bulky and brittle secondary phases, it is almost impossible to use the high-Zn high-Mg Al alloys in industries [8] .

To overcome the above disadvantages of Al-Zn-Mg-Cu alloys, nanoparticles have been introduced to modify (i.e nanotreat) the alloy systems very recently [9] , [10] . However, even with the addition of various nanoparticles, [9] , [11] successful manufacturing of bulk high-Zn high-Mg Al alloys has yet been demonstrated. More importantly, the systematic experimental analysis to link the phase evolution during the post-processing isn't available, which adds to the difficulty of developing high-Zn high-Mg light Al alloys with desired strength. [10] Due to

Al-Zn-Mg-Cu alloys (7xxx series aluminum alloys) are extensively studied and used for aeronautical applications where high strength and toughness are required [1] . These alloy systems are mainly composed of easily procured elements of Zn, Mg, and Cu, and the high strength with low cost for raw materials makes these alloys irreplaceable. It has been known that three ways are effective to optimize the properties of Al-Zn-Mg-Cu alloy: 1) Changing the main alloying elements of Zn, Mg, and Cu; 2) Tuning the trace elements (Sc, Cr, Zr, and Mn); 3) Minimizing the atomic-scale element impurity (Fe, Si) [2] . All these methods are important for achieving a suitable combination of precipitate formation like η(MgZn2), T(Al2Mg3Zn3), S(Al2CuMg) and θ(Al2Cu).

Recently, with strong needs for higher strength and balanced other properties, to mitigate the costs by the expensive trace elements, research on Al alloys with higher-Zn (>8.5 wt. %) and higher-Mg becomes a hot topic [3] - [6] . Certain theoretical calculations linking secondary phase show the promising future of using high-concentration Zn and Mg for Al alloy strengthening [7] . However, with more Zn and Mg into the aluminum alloy system, the alloy may be difficult to deform due to the brittle large precipitates. Without effective modifications for these bulky and brittle secondary phases, it is almost impossible to use the high-Zn high-Mg Al alloys in industries [8] .

To overcome the above disadvantages of Al-Zn-Mg-Cu alloys, nanoparticles have been introduced to modify (i.e nanotreat) the alloy systems very recently [9] , [10] . However, even with the addition of various nanoparticles, [9] , [11] successful manufacturing of bulk high-Zn high-Mg Al alloys has yet been demonstrated. More importantly, the systematic experimental analysis to link the phase evolution during the post-processing isn't available, which adds to the difficulty of developing high-Zn high-Mg light Al alloys with desired strength. [10] Due to

Al-Zn-Mg-Cu alloys (7xxx series aluminum alloys) are extensively studied and used for aeronautical applications where high strength and toughness are required [1] . These alloy systems are mainly composed of easily procured elements of Zn, Mg, and Cu, and the high strength with low cost for raw materials makes these alloys irreplaceable. It has been known that three ways are effective to optimize the properties of Al-Zn-Mg-Cu alloy: 1) Changing the main alloying elements of Zn, Mg, and Cu; 2) Tuning the trace elements (Sc, Cr, Zr, and Mn); 3) Minimizing the atomic-scale element impurity (Fe, Si) [2] . All these methods are important for achieving a suitable combination of precipitate formation like η(MgZn2), T(Al2Mg3Zn3), S(Al2CuMg) and θ(Al2Cu).

Recently, with strong needs for higher strength and balanced other properties, to mitigate the costs by the expensive trace elements, research on Al alloys with higher-Zn (>8.5 wt. %) and higher-Mg becomes a hot topic [3] - [6] . Certain theoretical calculations linking secondary phase show the promising future of using high-concentration Zn and Mg for Al alloy strengthening [7] . However, with more Zn and Mg into the aluminum alloy system, the alloy may be difficult to deform due to the brittle large precipitates. Without effective modifications for these bulky and brittle secondary phases, it is almost impossible to use the high-Zn high-Mg Al alloys in industries [8] .

To overcome the above disadvantages of Al-Zn-Mg-Cu alloys, nanoparticles have been introduced to modify (i.e nanotreat) the alloy systems very recently [9] , [10] . However, even with the addition of various nanoparticles, [9] , [11] successful manufacturing of bulk high-Zn high-Mg Al alloys has yet been demonstrated. More importantly, the systematic experimental analysis to link the phase evolution during the post-processing isn't available, which adds to the difficulty of developing high-Zn high-Mg light Al alloys with desired strength. [10] Due to

Al-Zn-Mg-Cu alloys (7xxx series aluminum alloys) are extensively studied and used for aeronautical applications where high strength and toughness are required [1] . These alloy systems are mainly composed of easily procured elements of Zn, Mg, and Cu, and the high strength with low cost for raw materials makes these alloys irreplaceable. It has been known that three ways are effective to optimize the properties of Al-Zn-Mg-Cu alloy: 1) Changing the main alloying elements of Zn, Mg, and Cu; 2) Tuning the trace elements (Sc, Cr, Zr, and Mn); 3) Minimizing the atomic-scale element impurity (Fe, Si) [2] . All these methods are important for achieving a suitable combination of precipitate formation like η(MgZn2), T(Al2Mg3Zn3), S(Al2CuMg) and θ(Al2Cu).

Recently, with strong needs for higher strength and balanced other properties, to mitigate the costs by the expensive trace elements, research on Al alloys with higher-Zn (>8.5 wt. %) and higher-Mg becomes a hot topic [3] - [6] . Certain theoretical calculations linking secondary phase show the promising future of using high-concentration Zn and Mg for Al alloy strengthening [7] . However, with more Zn and Mg into the aluminum alloy system, the alloy may be difficult to deform due to the brittle large precipitates. Without effective modifications for these bulky and brittle secondary phases, it is almost impossible to use the high-Zn high-Mg Al alloys in industries [8] .

To overcome the above disadvantages of Al-Zn-Mg-Cu alloys, nanoparticles have been introduced to modify (i.e nanotreat) the alloy systems very recently [9] , [10] . However, even with the addition of various nanoparticles, [9] , [11] successful manufacturing of bulk high-Zn high-Mg Al alloys has yet been demonstrated. More importantly, the systematic experimental analysis to link the phase evolution during the post-processing isn't available, which adds to the difficulty of developing high-Zn high-Mg light Al alloys with desired strength. [10] Due to 48th SME North American Manufacturing Research Conference, NAMRC 48 (Cancelled due to these drawbacks, current manufacturing methods including spray deposition [6] , powder metallurgy [12] , and above-room temperature post-processing [5] , [13] , [14] are hardly compatible with the eco-friendly large-scale fabrication needs. In addition, very limited mass-production methods (e.g. extrusion [15] ) are readily available for high-quality high-Zn high-Mg 7xxx alloy series.

Here, we prove the feasibility of manufacturing of Al-12Zn-3.7Mg-1Cu-0.56Mn-0.13Cr-0.14Zr (i.e. Al-12Zn-3.7Mg-1Cu) system. 12 wt.% Zn and 3.7 wt.% Mg are chosen for the effective secondary phase formation with η(MgZn2), as the prerequisite for high strength; Cu is lowered to 1 wt.% to allow the material naturally-aged and minimize S(Al2CuMg) phases. After hot rolling [16] or solution treatment (under temperatures of 460℃, 475℃, and 485℃) followed by natural aging, the microstructure and microhardness of the pure Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC are compared. The differential scanning calorimeter measurement (DSC) was carried out to analyze the solidification dynamics of the alloy, and the scanning electron microscopy (SEM) revealed the different microstructure evolution during these processes. The effective microhardness increase stands out with modified secondary phase in the Al-12Zn-3.7Mg-1Cu alloy with TiC nanoparticle incorporation. Further characterization by nanoindentation and focus ion beam (FIB) micro-pillar tests reveal the strength enhancement of Al-12Zn-3.7Mg-1Cu by 1.5 vol.% TiC nano-treating. [17] .

For Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC nanocomposites, the TiC nanoparticles are made in situ via the reaction of K2TiF6 and carbon black in molten aluminum at 800℃ for 2 hours. The Al-TiC master is cast out. Then, different element metals including Zn, Mg, Cu, Cr, Zr, and Mn are added to the master to obtain the Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC nanocomposites. Pure Al-12Zn-3.7Mg-1Cu is fabricated in the same fashion. The element composition is checked with inductively coupled plasma mass spectrometry (IC-PMS), and the results for Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC nanocomposites is shown in Table 1 .

After the alloying and casting, the materials will be hot rolled (with 60% thickness reduction) (i.e. "hot rolling-natural aging" or HR route) or solutionized (for 1-2h) (i.e. "solution treatment-natural aging" or ST route) under 460℃, 475℃, and 485℃, as shown in Fig. 1 . For HR, all the samples have the initial thickness of 1.0 ± 0.1 ; when the sample is rolled, the thickness reduces 0.1 each time, and the samples will be kept under the designed temperature for 5 before the next rolling cycle. The rolling speed is 7.5 / on MTI HR01 Hot Rolling Machine. From then on, all the samples are naturally aged for 20 days for the precipitate to form and evolve. Fig. 1 . The procedure illustration of solution treatment and hot rolling for Al-12Zn-3.7Mg-1Cu-0.56Mn-0.13Cr-0.14Zr alloy and its TiC nanocomposites

The microstructure of Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC nanocomposites immediately after solution treatment (ST) is summarized in Fig.2 . It's clear that under solution treatment at 460℃ (Fig. 2(a) ), 475℃ ( Fig. 2(b) ), and 485℃ (Fig. 2(c) ), the phase boundary of the Cu-and Zn-contained secondary phases is already blurred as the sample is being solutionized, and the size of these secondary phases are smaller when the solutionizing temperature goes higher. Importantly, Mg distributes reasonably uniformly after solution treatment, which is the prerequisite for the strengthening phase formation and effective processing during the latter natural aging. 

The microstructure after 20 days natural aging of Al-12Zn-3.7Mg-1Cu alloy and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC nanocomposites are shown in Fig. 3 .

For all the samples with TiC nanoparticles, the secondary phases show a smaller pattern as compared to the pure alloy, and this proves the effective phase control by the nanoparticles. [8] Consistently, during the hot rolling, the Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC nanocomposites delayed the rolling cracks to a larger strain and are more deformable by the rolling processing.

(Continued) Furthermore, by comparing the hot-rolled samples with the pure solution-treated ones, it shows that the precipitates in the hot rolled samples are smaller, more uniform in shape, and more homogeneously distributed. As a case study, Fig. 4 shows the precipitate distribution (by the precipitate size) of Al-12Zn-3.7Mg-1Cu/TiC nanocomposites by ST (Fig. 4(a) ) and HR ( Fig. 4(b) ) route under 475℃, respectively. It's clear that the HR route have the more uniform, finer precipitates centering the size of ≤ 1500 2 (over 70% of the precipitates, compared to 31% of those in pure Al-12Zn-3.7Mg-1Cu).

Especially for the hot-rolled samples under 460℃, the precipitate is invisible even under SEM, but its effects come out in microhardness as shown in Fig. 5(a) . This proves that the hot rolling can still be an effective processing method to optimize the properties of Al-12Zn-3.7Mg-1Cu alloy and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC nanocomposites. 

The microhardness evolution during the 20-day natural aging is recorded in Fig. 5 . For all temperatures from 460℃ to 485℃ (Fig. 5 (a)-(c) ), the natural aging is effective (compared with Zn-and Mg-low-content Al alloys) as the microhardness goes up, because the precipitate nucleation is enhanced via the heterogeneous interface at TiC nanoparticles (i.e. N) and the lowered free energy (i.e. ∆G) by the increased Zn and Mg chemical potential, according to Equation (1): Under the designed processing temperatures, the microhardness by "hot rolling-natural aging" is higher than purely by "solution treatment-natural aging", since hot rolling as the high-energy processing technique can introduce more defects synergetic with TiC nanoparticles and precipitates for strengthening.

The peak microhardness after natural aging is shown in Fig.  5(d) . Interestingly, for the solution-treated samples, the peak microhardness exists around 475℃, whereas the peak microhardness for hot-rolled samples comes at 460℃. This indicates the different contributions from deformation and temperature to the strengthening in Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC.

To fully demonstrate the feasibility of using the bulk manufacturing method to fabricate high-strength high-Zn high-Mg aluminum alloy, the mechanical properties including Young's modulus, Mohr's hardness, and compression strength.

The Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC processed via "ST route" at 475℃ (i.e. Solution treatment 475℃, followed by natural aging for 20 days) were chosen, because this processing route gives the peak-hardness state of this material and also introduces no work-hardening from rolling, as shown in Fig. 5(d) .

As shown in Fig. 6 , the Mohr's hardness is increased by 116.5% from 1.39 GPa to 3.01 GPa, and the Young's modulus is enhanced by 21.3% from 74.7 GPa to 90.6 GPa with the TiC nano-treating under the solution-treatment peak-hardness state. shows the typical micro-pillars fabricated by FIB for both Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC. The stress-strain curves and their post-test morphology have also been summarized in Fig. 7 (c)-(e). The yield compressive strength of Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC is much higher than that of Al-12Zn-3.7Mg-1Cu. As shown in Fig. 7 (d) and (e), the TiC nanoparticles could also effectively impede the slip plane movement (as marked by the less activated slip planes in yellow arrow). It can be clearly indicated that the nano-treating from TiC nanoparticles could effectively enhance the mechanical properties of the high-Zn high-Mg systems by this significant plan slip impediment. [18] More importantly, when compared with the compression strrengths from the similar systems in the previous studies [12] , the compression strength and strain in our systems are much better. The advantages by the synergy between highconcentration precipitates and TiC nanoparticles are clear.

(Continued) Fig. 7 . Focused ion beam-machined micro-pillar for (a) Al-12Zn-3.7Mg-1Cu and (b) Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC; The compression stress-strain curve for (c) Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol.5 TiC; High-resolution SEM images after compression test for (d) Al-12Zn-3.7Mg-1Cu and (e) Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC.

To better analyze the precipitation dynamics of the Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC systems, DSC measurement was carried out at 10℃/min quasiequilibrium, as shown in Fig. 8 . It's indicated that the secondary phase including aged precipitates has the dynamic melting window from ~480-495℃.

Recalling the peak microhardness in Fig. 5(d) , the solution treated samples have the highest microhardness with a processing temperature around 475℃, since this temperature is near the highest temperature to facilitate the solutionization procedure while not introducing direct secondary phase melting. [19] This state will be the highest entropy to decrease (c) the free energy. By Equation (1), this will result in a higher precipitate nucleation speed and higher microhardness. 

In this sense, the trend for microhardness after hot rolling can be expressed with:

Where v is the activation volume during hot rolling and mainly determined by the matrices. σ is the stress for deformation, and ε for the hot rolled strain (60% in our case). Hence, exp (∆ ⁄ ) denotes the strengthening effect from deformation for the residual secondary phase after solution treatment. Since 460℃ has the highest concentration of residual undissolved secondary phase (See Fig. 2(a) ), HV after hot rolling for Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC should be the most significant as shown in Fig. 5(d) .

When it comes to the enhanced mechanical properties of Al-12Zn-3.7Mg-1Cu by 1.5 vol.% TiC nano-treating, this synergy among the high-concentration secondary phases/precipitates (i.e. ∆G(T)), the strain after rolling (i.e. ∆U), and nano-treating from TiC (i.e. N) is also responsible. Therefore, as shown in Fig. 9 , with the incorporation of TiC and subsequent phase modification, nano-treating could utilize and unify the benefits from precipitate concentration and common processing techniques. Microscopically, the interaction of TiC nanoparticles with the defects and dislocations for traditional strengthening mechanisms (i.e. precipitate strengthening and work hardening) makes high-strength Al-Zn-Mg-Cu system possible. In brief, the road map shown in Fig. 9 with effective nano-treating is promising and feasible. Fig. 9 . Illustration of the processing flow and road map for propertyenhancement synergy in the studied Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol% TiC.

The new manufacturing route of "solution treatmentnatural aging" and "hot rolling-natural aging" for processing the high-Zn high-Mg Al-12Zn-3.7Mg-1Cu and Al-12Zn-3.7Mg-1Cu/1.5 vol.% TiC for higher strength is feasible; By controlling the secondary phase during processing, the addition of 1.5 vol.% TiC is validated to offer the alloy system overall enhanced mechanical properties. The synergy among the highconcentration secondary phases/precipitates, the strain after rolling, and nano-treating from TiC is responsible for the property improvement. In general, the nanoparticle incorporation for high strength Al alloys adds to another important design dimension (see Fig. 9 ) to the manufacturing space of novel high-Zn and high-Mg 7xxx series aluminum alloys.

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We sincerely thank Mr. Narayanan Murali at UCLA for his careful proofreading and useful comments for this manuscript. The raw materials and related fabrication equipment are provided by MetaLi LLC. This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors.