materials

Article Enhanced High-Temperature Mechanical Properties of Al–Cu–Li through T1 Coarsening Inhibition and Ce-Containing Intermetallic Refinement

Xinxiang Yu 1,2,* , Zhiguo Zhao 3, Dandan Shi 3, Han Dai 1,2, Jie Sun 1 and Xiaoyan Dong 1 1 Laboratory of Advanced Light Alloy Materials and Devices, Yantai Nanshan University, Longkou 265713, China; [email protected] (H.D.); [email protected] (J.S.); [email protected] (X.D.) 2 Postdoctoral Station of Nanshan Group Co., Ltd., Longkou 265706, China 3 Testing Center, Hang Xin Material Technology Co., Ltd., Longkou 264006, China; [email protected] (Z.Z.); [email protected] (D.S.) * Correspondence: [email protected]; Tel.: +86-0535-859-0938

 Received: 26 April 2019; Accepted: 5 May 2019; Published: 9 May 2019 

Abstract: The effects of the addition of 0.29 wt% Ce on the high-temperature mechanical properties of an Al–Cu–Li alloy were investigated. Ce addition contributes to T1 (Al2CuLi) phase coarsening inhibition and Ce-containing intermetallic refinement which greatly improved the thermal stability and high-temperature deformation uniformity of this alloy. On the one hand, small Ce in solid solution and segregation at phase interface can effectively prevent the diffusion and convergence of the main element Cu on T1 phase during thermal exposure. Therefore, the thermal stability of Ce-containing alloy substantiality improves during thermal exposure at the medium-high-temperature stage (170 ◦C to 270 ◦C). On the other hand, the increment of the tensile elongation in Ce-containing alloy is much greater than that in Ce-free alloy at high temperatures tensile test, because the refined Al8Cu4Ce intermetallic phase with high-temperature stability are mainly located in the fracture area with plastic fracture characteristics. This work provides a new method for enhancing high-temperature mechanical properties of Al–Cu–Li alloy which could be used as a construction material for high-temperature structural components.

Keywords: Al–Cu–Li alloy; cerium; thermal stability; intermetallic; microstructure

1. Introduction Al–Cu–Li alloy has a number of advantages over conventional aluminum alloy, such as lower , higher strength, higher fracture toughness, better fatigue, and corrosion resistance [1,2]. These properties contribute to the improvement of payload and fuel efficiency, thereby having made Al–Cu–Li alloy a competitive alternative to conventional 2xxx and 7xxx series materials in new aircraft designs [3]. The evaluation of thermal stability of Al–Cu–Li alloy is urgently needed for the potential application on fuselages, wing structures, bulkheads near the engine, as well as other high-temperature structural components which require high-temperature resistance in modern fighter aircraft [4]. So far the lower exposure temperatures in the range of 343 K to 358 K (70 ◦C to 85 ◦C) have been chosen to simulate the environment to which the wings and fuselage structures of commercial aircraft are exposed [5]. Reports have also been made on thermal exposure of Al–Cu–Li alloys at medium-high-temperatures (below 200 ◦C) [6,7]. However, materials near the engine, like bulkheads, are usually exposed to high-temperature environments (>200 ◦C). Therefore, the changes of microstructure and performance of Al–Cu–Li alloy in this temperature range requires further clarification, yet it has seldom been reported. The high-temperature (>70 ◦C) mechanical properties of Al–Cu–Li alloy depends on the thermal stability of its precipitates which includes T1 (Al2CuLi), δ0 (Al3Li), and θ0 (Al2Cu) [8]. Among them,

Materials 2019, 12, 1521; doi:10.3390/ma12091521 www.mdpi.com/journal/materials Materials 2019, 12, 1521 2 of 14 the T1 phase is the most stable [9], as the plates of T1 are very thin (~1.3 nm) and their thickness is very stable with aging at temperatures below ~170 ◦C[2]. Therefore, we chose T1 as the dominant strengthening phase in Al–Cu–Li alloy. The T1 phase can be promoted by increasing Cu/Li ratio [10] and small pre-stretching prior to artificial aging [2,11]. In addition, the thermal stability of T1, a Cu-containing precipitate, can be further improved by adding rare elements. A small addition of Sc [12], Y[13], La [14], and Er [15] can decrease the diffusion rate of the main alloying element Cu, thereby delaying the coarsening of Cu-containing precipitates in Al alloy. Ce as rare earth element also has the same function [16,17], so the coarsening process of the Cu-containing precipitate in Ce-containing alloy at high temperature can be delayed [18] and the thermal stability of these alloys can be improved, e.g., the addition of Ce inhibits the coarsening of Ω (Al2Cu) precipitate in Al–Cu–Mg–Ag alloy [19] and T1 (Al2CuLi) precipitate in Al–Cu–Li alloy [20]. Ce can also contribute to the refinement of intermetallic of aluminum alloys, e.g., a refined eutectic structure in 7055Al alloy [21], Al–6.7Zn–2.6Mg–2.6Cu (wt %) alloy [22], and Al–Cu–Li alloy [23] have been reported. A large number of small Al8Cu4Ce phases were also formed in Ce-containing Al–Cu–Li alloy after homogenization in our previous work [24]. These high melting point intermetallics do not coarsen evidently at elevated temperatures [25], which might strengthen grain boundaries and serve as obstacles to slip and creep at high temperatures to improve high-temperature mechanical properties of Ce-containing Al–Cu–Li alloy. Refined Ce-containing intermetallic improves the high-temperature mechanical properties of the Sn–3.8Ag–0.7Cu solder joint [26], but research on the Al–Cu–Li alloy from this approach has seldom been made. The purpose of this work is to investigate the effect of Ce addition on the inhibition of the main strength phase T1 coarsening during high-temperature exposure in high Cu/Li ratio Al–Cu–Li alloy at T8 temper (6% pre-stretching), and the effect of Ce on the diffusion of Cu in this alloy by a designed diffusion couple. Moreover, the enhanced deformation uniformity and strength of the Al–Cu–Li alloy owing to the refined Ce-containing intermetallic during high-temperature tensile were also studied in this paper.

2. Materials and Methods The materials used in this study were a 0.29 wt% Ce-containing alloy and a Ce-free Al–Cu–Li alloy, as shown in Table1. Master alloys of Al–Cu (50 wt%), Al–Zr (3.29 wt%), Al–Ce (10 wt%), and pure Ag, Mg, Li, and the remaining Al were melted in a vacuum induction melting furnace in a controlled Ar gas atmosphere, in a high-purity graphite crucible [1]. The casting was performed in Ar, using a Cu mold surrounded by cooling water [27]. The ingots were homogenized through a two-step homogenization course at 470 ◦C, 8 h and 510 ◦C, 16 h in a salt bath, followed by air cooling. After that, the ingots were hot-rolled from 23 mm to 4.5 mm and finally cold-rolled to 2.2 mm in thickness. Subsequently, the plates were solution-treated at 520 ◦C for 1 h, water-quenched, and finally aged at 150 ◦C with pre-stretching at 6% (T8 temper).

Table 1. Chemical composition of investigated alloys (wt%).

Alloys Li Cu Ag Mg Zr Ce Ce-free alloy 1.29 4.62 0.41 0.39 0.14 – Ce-containing alloy 1.31 4.63 0.38 0.40 0.13 0.29

In order to assess the thermal stability of the experimental alloys during thermal exposure, several specimens 10 mm 15 mm 2 mm in size, were cut from the plate at T8 temper. These specimens were × × then subjected to different over-aging heat treatments within the temperature range of 70 ◦C–320 ◦C, the soaking time ranging from 250 h up to 1300 h. The microhardness was measured by HV-1000 (Shanghai Yanrun Light Machinery Technology Co., Ltd., Shanghai, China), with a loading force of 0.2 kg and a loading time of 20 s. Tensile tests of specimens were performed at the temperatures from 25 ◦C to 300 ◦C, an Instron machine (INSTRON, Boston, MA, USA) equipped with incubator was used. The specimens were taken in the longitudinal direction from alloy plates at T8 temper. Materials 2019, 12, x FOR PEER REVIEW 3 of 14 with incubator was used. The specimens were taken in the longitudinal direction from alloy plates at Materials 2019, 12, 1521 3 of 14 T8 temper. The microstructure of the precipitate was examined by transmission electron microscopy (JEOL-2100F,The microstructure JEOL Corp., of Tokyo, the precipitate Japan) atwas an accelerating examined by voltage transmission of 200 kV. electron TEMmicroscopy foils were prepared(JEOL-2100F, by the JEOL twin–jet Corp., polishing Tokyo, Japan) technique at an using accelerating a solution voltage of 25% of 200nitric kV. acid TEM and foils 75% were methanol prepared at −by30 the°C, twin–jet70–80 mA, polishing and 15 techniqueV. Grain structure using a solution of the specimens of 25% nitric was aciddetermined and 75% after methanol being atanodized30 C, − ◦ with70–80 Barker’s mA, and reagent 15 V. Grain(1.8% structureHBF4 solution of the with specimens a voltage was ofdetermined 20–30 V and after a current being anodizedof 0.5–1.5 withmA, Quanzhou,Barker’s reagent China) (1.8% and HBF viewed4 solution with with polarized a voltage ligh oft 20–30 on optical V and amicroscope current of 0.5–1.5 (ZEISS, mA, Oberkochen, Quanzhou, Germany).China) and The viewed evaluation with polarized of the fracture light on features optical microscope of the high-temperature (ZEISS, Oberkochen, tensile Germany).specimen was The performedevaluation ofwith the fracturean FEI-Nova features Nano of the SEM high-temperature 450 scanning tensile electron specimen microscope was performed (FEI Company, with an Hillsboro,FEI-Nova NanoOR, USA). SEM 450 scanning electron microscope (FEI Company, Hillsboro, OR, USA). To verifyverify thethe inhibitinginhibiting e ffeffectect of of small small Ce Ce addition addition on on T1 T1 phase phase coarsening, coarsening, a di aff usiondiffusion couple couple was wasdesigned. designed. Firstly, Firstly, a pure a aluminum pure aluminum plate was plat sandwichede was sandwiched between a Ce-freebetween alloy a Ce-free and a Ce-containing alloy and a Ce-containingalloy (The oxide alloy layer (The at theoxide joint layer of alloyat the sheet joint wasof alloy polished sheet cleanwas polished by water clean grinding by water sand). grinding Then it sand).was drilled Thenand it was fixed drilled with copperand fixed wire. with After short-time wire. preheating,After short-time it was preheating, tightly welded it was by tightly rolling weldedand then by annealed rolling and at 520 then◦C forannealed 6 h. The at formation520 °C for of 6 ah. new The oxide formation layer wasof a prevented new oxide by layer inert was gas preventedprotection duringby inert annealing gas protection and rolling. during The annealin cross sectiong and ofrolling. the specimen The cross was section polished of the in accordance specimen waswith polished metallographic in accordance standard with and metallographic then observed standard by SEM. and then observed by SEM.

3. Results Results and and Discussion Discussion The initial microstructures of Ce-containing and Ce Ce-free-free alloys at T8 temper were characterized both by TEMTEM andand three-dimensionalthree-dimensional OM OM investigations, investigations, as as shown shown in in Figure Figure1. Figure 1. Figure1a shows 1a shows that thethat themain main precipitate precipitate in two in alloys two isalloys uniformly is uniformly distributed distributed T1 phase atT1 T8 phase temper. at AT8 small temper. pre-stretching A small pre-stretchingbefore aging induces before dislocation aging induces network dislocation within the network matrix, which within acts the as heterogeneousmatrix, which nucleation acts as heterogeneoussites for the T1 nucleation phase and greatlysites for promotes the T1 phase T1 phase and greatly precipitation promotes [28]. T1 phase precipitation [28].

Figure 1.1. MicrostructuresMicrostructures comparison comparison of of Ce-containing Ce-containing and and Ce-free Ce-free alloys alloys at T8 at temper: T8 temper: (a) TEM (a) imageTEM imageof T1 phase; of T1 phase; (b) three (b) direction three direction optical optical micrographs micrographs of grain of structure. grain structure.

Although the distribution characteristics characteristics of of T1 T1 phase phase are are not affected affected by Ce addition, many fiber-likefiber-like unrecrystallized microstructures are are foun foundd distributed along along the the rolling direction in the Ce-containing alloyalloy whilewhile thethe recrystallized recrystallized grains grains are are primarily primarily observed observed in Ce-freein Ce-free alloy alloy (Figure (Figure1b), 1b),Clearly Clearly Ce addition Ce addition effectively effectively improves improves the recrystallization the recrystallization resistance resistance of Al–Cu–Li of Al–Cu–Li alloys duringalloys duringsolid solution solid solution treatment treatment [1]. Therefore, [1]. Therefore, the initial the microhardnessinitial microhardness of Ce-containing of Ce-containing alloy (189 alloy HV (1890.2) HVis greater0.2) is greater than that than of that Ce-free of Ce-free alloy (177alloy HV (1770.2 ).HV Moreover,0.2). Moreover, this is this mainly is mainly attributed attributed to the to grain the grain size sizedifference difference between between them. them.

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The thermal stability curves of Ce-containing and Ce-free alloys at T8 temper which are characterized by by microhardness microhardness and and obtained obtained in in the the range range 70–320 70–320 °C◦ Cfor for soaking soaking time time from from 250 250 h to h 1300to 1300 h, h,are are reported reported in inFigure Figure 2.2 .The The microhardness microhardness of of Ce-containing Ce-containing alloy alloy is found to bebe muchmuch higher thanthan that that of of Ce-free Ce-free alloy alloy within within the temperature the temperature range from range 170 from◦C to 270170 ◦°CC. Atto temperatures270 °C. At temperaturesno higher than no 120 higher◦C, good than thermal 120 °C, stability good therma is observedl stability in both is alloys.observed As thein both temperature alloys. furtherAs the temperatureincreases to 320 further◦C,the increases hardness to 320 values °C, ofthe the hardness two alloys values become of the similar two alloys again. become The thermal similar stability again. Theof the thermal Al–Cu–Li stability alloy of is the substantially Al–Cu–Li alloy improved is substantially by small Ce improved addition by during small thermalCe addition exposure during at thermalmedium-high-temperature exposure at medium-high-temperature stage (170 ◦C to 270 ◦stageC). (170 °C to 270 °C).

Figure 2. Thermal stabilitystability curvescurves of of Ce-containing Ce-containing and and Ce-free Ce-free alloys alloys during during thermal thermal exposure exposure at the at thetemperature temperature range range 70–320 70–320◦C for°C soakingfor soaking time time from from 250 250 h to h 1300 to 1300 h. h. The precipitate phase characteristics of Ce-containing and Ce-free alloys after thermal exposure The precipitate phase characteristics of Ce-containing and Ce-free alloys after thermal exposure at 170 C for 1000 h and 220 C for 250 h were investigated respectively by using TEM, as shown in at 170 °C◦ for 1000 h and 220 °C◦ for 250 h were investigated respectively by using TEM, as shown in Figure3. The thickness of the precipitates in Ce-containing alloy is 25.40 6.52 nm (average standard Figure 3. The thickness of the precipitates in Ce-containing alloy is± 25.40 ± 6.52 nm (average± ± deviation), which is much lower than that in Ce-free alloy (31.22 10 nm) after exposure at 170 C for standard deviation), which is much lower than that in Ce-free alloy± (31.22 ± 10 nm) after exposure◦ at 1000 h. Although the diameter (length) of the precipitates in Ce-free alloy (114.45 29.70 nm) is slightly 170 °C for 1000 h. Although the diameter (length) of the precipitates in Ce-free± alloy (114.45 ± 29.70 more than that in Ce-containing alloy (106.35 nm 28.19 nm), the effect of thickness increment of the nm) is slightly more than that in Ce-containing alloy± (106.35 nm ± 28.19 nm), the effect of thickness precipitates (22.91%) is much more potent compared with that of their diameter increment (7.62%). increment of the precipitates (22.91%) is much more potent compared with that of their diameter After exposure at 220 C for 250 h, the thickness of precipitates in Ce-containing alloy becomes 16.85 increment (7.62%). After◦ exposure at 220 °C for 250 h, the thickness of precipitates in Ce-containing± 3.41 nm, which is much lower than that in Ce-free alloy (27.58 7.25 nm). Moreover, the diameter alloy becomes 16.85 ± 3.41 nm, which is much lower than that± in Ce-free alloy (27.58 ± 7.25 nm). of precipitates in Ce-containing alloy turns to 196.11 55.65 nm, which is much greater than that in Moreover, the diameter of precipitates in Ce-containing± alloy turns to 196.11 ± 55.65 nm, which is Ce-free alloy (113.46 35.96 nm). Ce addition is beneficial to slowing down the coarsening rate of much greater than that± in Ce-free alloy (113.46 ± 35.96 nm). Ce addition is beneficial to slowing down the main phase T1 in the Al–Cu–Li alloy. These plate phase T1 precipitate along the {111} and it is the coarsening rate of the main phase T1 in the Al–Cu–Li alloy. These plate phase T1 precipitateα non-deformable [29], the “strengthening stress τp” produced by these phases can be expressed as [30]. along the {111}α and it is non-deformable [29], the “strengthening stress τp” produced by these phases can be expressed as [30].  1 1 3  b 2 2 0.079D τp = 0.12G f + 0.70(D/t) 2 fν + 0.12(D/t) f ln (1) 1/2 ν 1 1 3 ν bD220.079 r τ =++0.12(GDtfDtDt) 0.70( / )2 0.12( / ) ln 0 p 1/2 ffννν (1) ()Dt r0 D and t refer to the diameter and thickness of T1 phase respectively. Formula (1) shows that the strengthening stress varies inversely with the thickness and directly with the diameter of the T1 phase. Because Ce addition is beneficial to slowing down the coarsening rate of the T1 phase in the Al–Cu–Li alloy, the thickness of the precipitates of T1 in the Ce-containing alloy is much lower than that in Ce-free alloy during thermal exposure at 170 ◦C for 1000 h and 220 ◦C for 250 h. Since Cu is one of the main elements in the T1 phase, the coarsening rate of the main strengthening phase T1 in the Al–Cu–Li

Materials 2019, 12, 1521 5 of 14 alloy can be reduced through inhibiting the diffusion and convergence of the main element Cu on T1 phase during medium-high temperature thermal exposure. To prove this inference, a diffusion couple was designed to evaluate the effect of Ce addition on the diffusion of Cu in the Al–Cu–Li alloys, as shown in Figure4. Two di ffusion interfaces were taken as the starting point, and a quantitative EDS (Energy scattering X-ray spectrum) analysis was done along the solid line arrow in every 10 µm in Figure4. The di ffusion rate of Cu atoms from Ce-containing alloy to pure Al is much lower than that from Ce-free alloy to pure Al, as shown in the mapping of Cu distribution and Cu content distribution curve along the solid line arrow. So the reduced diffusion rate of Cu in the Al–Cu–Li alloy through Ce Materialsaddition 2019 was, 12, confirmedx FOR PEER byREVIEW the di ffusion couple experiment. 5 of 14

FigureFigure 3. 3. TheThe precipitate precipitate phase phase coarsening coarsening characteristics characteristics of of (a (a) )Ce-containing Ce-containing and and (b (b) )Ce-free Ce-free alloys alloys

afterafter exposure exposure at at 170 170 °C◦C for for 1000 1000 h h and and 250 250 °C◦C for for 250 250 h. h.

DThe and eff tect refer of Ceto the addition diameter is attributed and thickness to the of strong T1 phase interaction respectively. energy Formula between (1) rare shows earth that Ce theand strengthening vacancies in stress solid solutionvaries inversely matrix, with which the can thickness prevent and the directly migration with of the dislocation diameter and of the solute T1 phase.atom [Because16,17]. HenceCe addition small is Ce beneficial in solid to solution slowing can down prevent the coarsening the diffusion rate of of major the T1 element phase in Cu the in Al–Cu–Liexperimental alloy, alloys, the thickness and ultimately of the inhibitprecipitates the coarsening of T1 in the of strengtheningCe-containing phasesalloy is T1 much in the lower Al–Cu–Li than thatalloy. in Furthermore,Ce-free alloy during due to thethermal large exposure size mismatch at 170 of °C Ce for and 1000 Al h atoms, and 220 Ce °C tends for 250 to segregate h. Since Cu at the is onecrystallographic of the main elements interfaces in forthe relievedT1 phase, lattice the coar strain.sening The rate segregation of the main of strengthening Ce at the edge phase interface T1 in of theprecipitates Al–Cu–Li would alloy ecanffectively be reduced restrict through its lateral inhibi growthting and the inhibit diffusion its coarsening. and convergence This would of the likewise main elementdecrease Cu the on copper T1 phase diffusion during rates medium-high as it decreases temperature the tendency thermal for higher exposure. diffusive To flux prove between this inference,precipitates a diffusion looking to couple coarsen. was Therefore, designed the toimproved evaluate the thermal effect stability of Ce addition of the Al–Cu–Li on the diffusion alloy during of Cuthermal in the exposure Al–Cu–Li at alloys, the medium-high-temperature as shown in Figure 4. Two stage diffusion is attributed interfaces to the were coarsening taken as inhibitionthe starting of point,the main and strengthening a quantitative phases EDS (Energy T1 by reducing scattering the X-ray diffusion spectrum) rate of analysis Cu element was anddone restricting along the lateralsolid linegrowth arrow through in every small 10 Ceμm addition. in Figure 4. The diffusion rate of Cu atoms from Ce-containing alloy to pure Al is much lower than that from Ce-free alloy to pure Al, as shown in the mapping of Cu distribution and Cu content distribution curve along the solid line arrow. So the reduced diffusion rate of Cu in the Al–Cu–Li alloy through Ce addition was confirmed by the diffusion couple experiment. The effect of Ce addition is attributed to the strong interaction energy between rare earth Ce and vacancies in solid solution matrix, which can prevent the migration of dislocation and solute atom [16,17]. Hence small Ce in solid solution can prevent the diffusion of major element Cu in experimental alloys, and ultimately inhibit the coarsening of strengthening phases T1 in the Al–Cu– Li alloy. Furthermore, due to the large size mismatch of Ce and Al atoms, Ce tends to segregate at the crystallographic interfaces for relieved lattice strain. The segregation of Ce at the edge interface of precipitates would effectively restrict its lateral growth and inhibit its coarsening. This would likewise decrease the copper diffusion rates as it decreases the tendency for higher diffusive flux between precipitates looking to coarsen. Therefore, the improved thermal stability of the Al–Cu–Li alloy during thermal exposure at the medium-high-temperature stage is attributed to the

Materials 2019, 12, x FOR PEER REVIEW 6 of 14

Materialscoarsening2019 , inhibition12, 1521 of the main strengthening phases T1 by reducing the diffusion rate of6 ofCu 14 element and restricting lateral growth through small Ce addition.

Figure 4. (a) Diagram of diffuse couple model; (b) mapping of Cu distribution and (c) Cu content Figure 4. (a) Diagram of diffuse couple model; (b) mapping of Cu distribution and (c) Cu content distribution curve along the solid line arrow in pure Al near the edge of the Ce-containing alloy and distribution curve along the solid line arrow in pure Al near the edge of the Ce-containing alloy and the Ce-free alloy respectively. the Ce-free alloy respectively. Figure5 is the microstructure and element mapping of precipitates in Ce-containing and Ce-free Figure 5 is the microstructure and element mapping of precipitates in Ce-containing and alloys after exposed at 270 ◦C for 250 h. Although the main strength phase T1 has been completely dissolved,Ce-free alloys a few after new exposed particles at containing 270 °C for Al, 250 Cu, h. andAlthough Ce with the an main average strength size of phase about T1 80 has nm werebeen foundcompletely dispersed dissolved, in the matrix.a few new Moreover, particles a large containing number Al, of smallerCu, and and Ce uniformlywith an average distributed size particlesof about in80 thenm matrixwere found were alsodispersed observed in the in Ce-containingmatrix. Moreover, alloy, a as large shown number in Figure of smaller5a. In contrastand uniformly to the casedistributed for Ce-containing particles in alloy,the matrix some coarsewere also rod-like observ phasesed in Ce-containing with “fingerprint alloy, patterns” as shown can in still Figure be seen 5a. inIn Ce-freecontrast alloy, to the which case are for dissolution Ce-containing marks alloy, of the some residual coarse phases rod-like of severely phases coarsened with “fingerprint T1 phase. Onlypatterns” very can few still coarse be AlCuZrseen in particlesCe-free alloy, with thewhich size are of 200dissolution nm, and somemarks small of the unevenly residual distributed phases of particlesseverely cancoarsened be seen T1 in Ce-freephase. Only alloy, very as shown few coarse in Figure AlCuZr5b. The particles hardness with di ff theerence size between of 200 nm, the twoand some small unevenly distributed particles can be seen in Ce-free alloy, as shown in Figure 5b. The alloys is about 15 HV0.2 (strength value is about 50 MPa), which is basically due to the contribution of thehardness newly difference formed AlCuCe between phase the in two Ce-containing alloys is about alloy. 15 Since HV the0.2 (strength AlCuCe particles value is and about those 50 smaller MPa), newlywhich is observed basically particles due to the are contribution evenly dispersed of the innewly the alloyformed and AlCuCe remain phase stable in in Ce-containing high temperature, alloy. Ce-containingSince the AlCuCe alloy particles still exhibits and those better smaller thermal newl stabilityy observed than Ce-free particles alloy are after evenly thermal dispersed exposure in the at alloy and remain stable in high temperature, Ce-containing alloy still exhibits better thermal 270 ◦C for 250 h. stability than Ce-free alloy after thermal exposure at 270 °C for 250 h.

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FigureFigureFigure 5.5. TheThe characteristicscharacteristics and andand STEM-EDS STEM-EDSSTEM-EDS mapping mappimappingng analysis analysisanalysis of ofprecipitatesof precipitatesprecipitates in ( inain) Ce-containing((aa)) Ce-containingCe-containing and and (b) Ce-free alloys after exposure at 270 °C for 250 h. and(b) Ce-free(b) Ce-free alloys alloys after after exposure exposure at 270 at ◦270C for °C250 for h.250 h.

InInIn addition, addition,addition, the thethe two twotwo experimental experimentalexperimental alloys alloysalloys exhibit exhibitexhibit excellent excellentexcellent thermal thermalthermal stability stabilitystability during duringduring exposure exposureexposure at atat temperaturestemperaturestemperatures belowbelowbelow 120 120120◦C, °C,°C, as shownasas shownshown in Figure inin FigureFigure2. So the 2.2. microstructuresSoSo thethe microstructuremicrostructure of the alloysss ofof afterthethe alloys exposurealloys afterafter at exposure120exposure◦C for at 1300at 120120 h were°C°C forfor also 13001300 observed hh werewere by also TEM,also observobserv as showneded byby in FigureTEM,TEM, 6 as.as The shownshown stable inin strengthening FigureFigure 6.6. TheThe phase stablestable T1 strengtheningofstrengthening the Al–Cu–Li phasephase alloy T1T1 does ofof thethe not Al–Cu–LiAl–Cu–Li coarsen in alloyalloy both doesdoes experimental notnot coarsencoarsen alloys inin bothboth and experimental aexperimental large number alloysalloys of densely andand aa largedistributedlarge numbernumber particles ofof denselydensely appear distributeddistributed in the matrix, particlesparticles this appearappear is the main inin thethe reason matrix,matrix, why thisthis the isis the thermalthe mainmain stability reasonreason whywhy of these thethe ° thermalalloysthermal is stability maintainedstability of of these these during alloys alloys exposure is is maintained maintained at 120 ◦C. during during The relative exposure exposure Cu content at at 120 120 ° ofC.C. the The The newly relative relative formed Cu Cu content content particle of of is thecomparedthe newlynewly formed semi-quantitativelyformed particleparticle isis comparedcompared with that semi-quantit ofsemi-quantit the matrixativelyatively and plate withwith phase thatthat T1ofof the bythe STEM-EDS, matrixmatrix andand asplateplate shown phasephase in T1FigureT1 by by STEM-EDS, STEM-EDS,7. The Cu content as as shown shown of thisin in Figure Figure new particle 7. 7. The The Cu isCu higher content content than of of this this that new new of T1 particle particle phase, is is thus higher higher the than particlethan that that might of of T1 T1 phase, thus the particle might be the square-shaped Al5Cu6Li2 phase [31]. The Al–Cu–Li alloys phase,be the square-shapedthus the particle Al 5miCught6Li 2bephase the [square-shaped31]. The Al–Cu–Li Al5Cu alloys6Li2 phase exhibit [31]. excellent The Al–Cu–Li thermal stability alloys ° exhibitduringexhibit excellent exposureexcellent thermal atthermal temperatures stability stability belowduring during 120 exposure exposure◦C because at at temperatures temperatures there is no obvious below below 120 coarsening120 ° CC because because of the there there stable is is no T1no obvious coarsening of the stable T1 phase and the fine Al5Cu6Li2 phase is formed at the same time. obviousphase and coarsening the fine Al of5 Cuthe6 stableLi2 phase T1 phase is formed and atthe the fine same Al5Cu time.6Li2 phase is formed at the same time.

Figure 6. The precipitate characteristics of (a) Ce-containing and (b) Ce-free alloys after exposure at FigureFigure 6. TheThe precipitate precipitate characteristics characteristics of of ( (aa)) Ce-containing Ce-containing and and ( (bb)) Ce-free Ce-free alloys alloys after after exposure exposure at at 120° °C for 1300 h. 120120 ◦CC for for 1300 1300 h. h.

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°° FigureFigure 7. 7. ((a(aa)) ) TEMTEM TEM imageimage image ofof of thethe the precipitatesprecipitates precipitates inin in Ce-freeCe-free Ce-free alloyalloy afterafter exposureexposure atat at 120120 120 ◦CCC forfor for 13001300 1300 h;h; h; STEM-EDSSTEM-EDS analysisanalysis ofofof the thethe matrix matrixmatrix ( b ((b),b),), cubic cubiccubic phase phasephase (c ),((cc and),), andand plate plateplate phase phasephase (d) ( as(dd)) point asas pointpoint 1, point 1,1, pointpoint 2, and 2,2, point andand pointpoint3, respectively. 3,3, respectively.respectively.

TheThe sizesizesize and andand morphology morphologymorphology of of intermetallicsof intermetallicsintermetallics in Ce-containing inin Ce-containingCe-containing and Ce-free andand Ce-freeCe-free alloys afteralloysalloys solid afterafter solution solidsolid solutionweresolution investigated werewere investigatedinvestigated by SEM and byby EDS,SEMSEM asandand shown EDS,EDS, in asasFigure showshown8n. inin The FigureFigure average 8.8. TheThe transverse averageaverage size transversetransverse of intermetallic sizesize ofof intermetallicparticlesintermetallic AlCuCe particlesparticles in Ce-containing AlCuCeAlCuCe inin alloy Ce-containingCe-containing is less than alloy 6alloyµm. isis However, lessless thanthan a few66 μμm. coarsem. However,However, intermetallic aa fewfew particles coarsecoarse intermetallicAlCuZrintermetallic about particlesparticles 25 µm are AlCuZrAlCuZr also observed aboutabout 2525 besides μμmm areare the alsoalso normal observedobserved small besidesbesides particles thethe in normalnormal Ce-free smallsmall alloy. particlesparticles Obviously, inin Ce-freeCe-freethe formation alloy.alloy. Obviously,Obviously, of coarse AlCuZr thethe formationformation intermetallic ofof coarsecoarse compounds AlCuZrAlCuZr intermetallicintermetallic was interrupted compoundscompounds by the strongwaswas interruptedinterrupted attraction bybybetween thethe strongstrong Cu and attractionattraction Ce atoms inbetweenbetween Ce-containing CuCu andand alloy CeCe [32atomsatoms], for theinin electronegativeCe-containingCe-containing alloyalloy differences [32],[32], betweenforfor thethe electronegativeelectronegativeCe and Cu are moredifferencesdifferences than that betweenbetween between CeCe Zrandand and CuCu Cu araree [ 17 moremore]. In thanthan addition, thatthat Cebetweenbetween and Zr ZrZr have andand the CuCu opposite [17].[17]. InIn addition,addition,microsegregation CeCe andand ZrZr patterns havehave thethe similar oppositeopposite to that microsegregatimicrosegregati of Mn and Zronon in patterns Alpatterns alloy duringsimilarsimilar solidification toto thatthat ofof MnMn [ 33andand], so ZrZr the inin Al CeAl alloyalloyand Zr duringduring atoms solidificationsolidification combine with [33],[33], Cu respectively soso thethe CeCe andand and ZrZr transport atomsatoms combinecombine in the opposite withwith CuCu direction respectivelyrespectively during andand solidification, transporttransport ininthereby thethe oppositeopposite avoiding directiondirection the formation duringduring of solidification,solidification, coarse AlCuZr therebythereby phase. Moreover,avoidingavoiding thethe the formation Ceformation addition ofof would coarsecoarse a ffAlCuZrAlCuZrect and phase.phase.refine the Moreover,Moreover, formation thethe of interdendritic CeCe additionaddition compounds wouldwould affectaffect due toandand supersaturated refinerefine thethe Ceformationformation atoms being ofof interdendriticinterdendritic expelled from compoundscompoundsthe solidified duedue grains toto andsupersaturatedsupersaturated accumulated CeCe at theatomsatoms front beingbeing of the expelledexpelled interface fromfrom between thethe solidsolidifiedsolidified and liquid grainsgrains during andand accumulatedaccumulatedsolidification atat [16 thethe]. Hence, frontfront ofof the thethe formation interfaceinterface and betweenbetween refinement solidsolid andand of the liquidliquid AlCuCe duringduring phases solidificationsolidification are promoted [16].[16]. through Hence,Hence, thetheCe addition.formationformation andand refinementrefinement ofof thethe AlCuCeAlCuCe phasesphases areare promotedpromoted throughthrough CeCe addition.addition.

Figure 8. Cont.

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Figure 8. The characteristics and EDS analysis of the intermetallic in (a) Ce-containing and (b) Ce-free alloys after solid solution at 520 °C for 1 h. Figure 8.8. TheThe characteristicscharacteristics and and EDS EDS analysis analys ofis theof intermetallicthe intermetallic in (a )in Ce-containing (a) Ce-containing and (b )and Ce-free (b) alloys after solid solution at 520 C for 1 h. ToCe-free study alloys the after effects solid of solution the refined◦ at 520 °CCe-containing for 1 h. intermetallic on high-temperature tensile properties of the experimental alloys at T8 temper, the tensile test curves of Ce-containing and To study the eeffectsffects ofof thethe refinedrefined Ce-containingCe-containing intermetallicintermetallic on high-temperature tensile Ce-free alloys were investigated at 25 °C, 100 °C, 175 °C, 250 °C, and 325 °C, respectively, as shown properties ofof thethe experimental experimental alloys alloys at T8at temper,T8 temp theer, tensilethe tensile test curves test curves of Ce-containing of Ce-containing and Ce-free and in Figure 9. The strength value of Ce-containing alloy is always higher than that of Ce-free alloy Ce-freealloys were alloys investigated were investigated at 25 ◦C, at 100 25◦ C,°C, 175 100◦C, °C, 250 175◦C, °C, and 250 325 °C,◦ C,and respectively, 325 °C, respectively, as shown in as Figure shown9. within the entire tensile test temperature range. The yield strength of the Ce-containing alloy is inThe Figure strength 9. The value strength of Ce-containing value of Ce-containing alloy is always alloy higher is thanalways that higher of Ce-free than alloy that withinof Ce-free the entirealloy higher than that of Ce-free alloy, by 21.5 MPa at 100 °C, by 27.5 MPa at 175 °C, by 13.5 MPa at withintensile testthe temperatureentire tensile range.test temperature The yield strengthrange. The of the yield Ce-containing strength of alloythe Ce-containing is higher than alloy that ofis 250 °C, and by 4.5 MPa at 325 °C. As the temperature increases, the strength difference between the higherCe-free than alloy, that by 21.5of Ce-free MPa at alloy, 100 ◦C, by by 21.5 27.5 MPa MPa at at 100 175 ◦°C,C, byby 13.527.5 MPaMPa at at 250 175◦C, °C, and by by 13.5 4.5 MPa MPa at two alloys decreases, whereas the elongation of Ce-containing alloy increased greatly from 7.5 to 250325 ◦°C,C. Asand the by temperature 4.5 MPa at 325 increases, °C. As the the strength temperatur differencee increases, between the thestrength two alloys difference decreases, between whereas the 17.5% at the testing temperature range between room temperature and 325 °C, while that of Ce-free twothe elongationalloys decreases, of Ce-containing whereas the alloy elongation increased of greatly Ce-containing from 7.5% alloy to 17.5% increased at the testinggreatly temperaturefrom 7.5 to alloy only increased from 7 to 13%. Ce addition effectively enhances the high-temperature ductility 17.5%range betweenat the testing room temperature temperature rang ande 325between◦C, while room that temperature of Ce-free and alloy 325 only °C, increasedwhile that from of Ce-free 7% to of the Al–Cu–Li alloy, and the strength of this alloy is properly maintained to a certain extent at the 13%.alloy Ceonly addition increased eff ectivelyfrom 7 to enhances 13%. Ce theaddition high-temperature effectively enhances ductility the of thehigh-temperature Al–Cu–Li alloy, ductility and the same time. ofstrength the Al–Cu–Li of this alloy alloy, is and properly the strength maintained of this to alloy a certain is properly extent at maintained the same time. to a certain extent at the same time.

Figure 9. High-temperature tensile curves of Ce-containing and Ce-free alloys at T8 temper. Figure 9. High-temperature tensile curves of Ce-containing and Ce-free alloys at T8 temper.

Figure 9. High-temperature tensile curves of Ce-containing and Ce-free alloys at T8 temper.

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To understand the effects of the refined Cu-containing intermetallic particles on the fracture modeTo of understandCe-containing the and effects Ce-free of the alloys refined after Cu-containing high temperature intermetallic tensile, particles the fracture on the morphologies fracture mode ofof the Ce-containing alloys were and investigated Ce-free alloys comparatively after high temperatureby secondary tensile, electron the fracture(SE) image morphologies and backscatter of the electronalloys were (BSE) investigated image, as shown comparatively in Figure by secondary10; Figure electron 11. The (SE) cleavage image facets, and backscatter which indicate electron a brittle(BSE) image,fracture, as are shown dominant in Figures in Ce-free 10 and alloy, 11. The while cleavage only a facets, few facets which are indicate seen in a brittleCe-containing fracture, alloyare dominant at the lower in Ce-free tensile alloy,test temperature. while only a Moreov few facetser, the are fine seen AlCuCe in Ce-containing particles alloyare located at the in lower an areatensile with test obvious temperature. plastic Moreover, fracture characteristics. the fine AlCuCe As particles the temperature are located reaches in an area 250 with°C, obviousa large numberplastic fractureof dimples characteristics. are observed As in the Ce-containing temperature reachesalloy. The 250 dimples◦C, a large in the number zone ofwith dimples densely are distributedobserved in small Ce-containing AlCuCe particles alloy. The decrease dimples in insize the and zone increase with densely in quantity, distributed which small is consistent AlCuCe withparticles the obvious decrease increase in size and of the increase number in quantity,of dimples which in tensile is consistent fracture with surface the obvious of A357 increase alloy by of additionthe number of rare of dimples earth element in tensile [34]. fracture As the surface temper ofature A357 rises alloy to by325 addition °C, the ofdimples rare earth further element increase [34]. inAs size the and temperature are distributed rises tomore 325 uniformly.◦C, the dimples The Al furtherCuCe particles increase become in size finer and areand distributedmore dispersive more becauseuniformly. the Thefine AlCuCe AlCuCe particles particles become are gradually finer and cr moreushed dispersive due to the because further the softening fine AlCuCe of the particles alloy matrixare gradually at higher crushed temperature due to the [1,35]. further Compared softening with of the Ce-containing alloy matrix atalloy, higher the temperature coarse AlCuZr [1,35]. particlesCompared are with prone Ce-containing to brittle fracture alloy, and the coarsethus woul AlCuZrd not particles contribute are to prone the plastic to brittle deformation fracture and of thusthe alloywould at not the contribute whole high-temperature to the plastic deformation tensile test of the. Therefore, alloy at the the whole small high-temperature high-temperature tensile stable test. AlCuCeTherefore, phase the smallin the high-temperatureCe-containing alloy stable is beneficial AlCuCe phaseto the inimprovement the Ce-containing of the alloyductility is beneficial of the Al– to Cu–Lithe improvement alloy in high-temperature of the ductility tensile of the Al–Cu–Litest. alloy in high-temperature tensile test.

Figure 10. Cont.

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FigureFigure 10. 10. FractureFracture morphology morphology of of ( (aa)) Ce-containing Ce-containing and and ( (bb)) Ce-free Ce-free alloys alloys after after high-temperature high-temperature

tensiletensile at at 100 100 °C,◦C, 175 175 °C,◦C, 250 250 °C,◦C, and and 325 325 °C.◦C.

InIn general,general, the the deformation deformation resistance resistance of metals of metals decreases decreases with temperature. with temperature. As the temperature As the temperatureincreases, the increases, cross-slip the and cross-slip climb of dislocationsand climb of become dislocations easier become and the easier bonding and force the betweenbonding atomsforce betweenis further atoms weakened, is further which weakened, lead to the which thermal lead activation to the thermal of alloy activation increase of [36 alloy]. A largeincrease number [36]. A of largestudies number [37–40 ]of on thestudies deformation [37–40] behavioron the deformati of aluminumon alloysbehavior at elevated of aluminum temperatures alloys indicated at elevated that temperaturesintermetallic phases indicated with that high-temperature intermetallic stabilityphases havewith ahigh-temperature strong pining effect stability on grain boundarieshave a strong and piningsub-grain effect boundaries, on grain whichboundaries can compensate and sub-grai forn the boundaries, grain boundary which weakening can compensate at high for temperatures. the grain boundaryA large number weakening of fine at rare high earth temperatures. intermetallic A phases large withnumber high of temperature fine rare earth stability intermetallic lead to improved phases withhigh high temperature temperature tensile stability properties lead of to the improved Al alloy, high e.g., temperature the tensile strength tensile ofproperties 2519 alloy of withthe Al 0.2 alloy, wt% e.g.,Y increased the tensile by 30% strength at 300 ◦ofC, 2519 which alloy was with attributed 0.2wt% to theY increased Al6Cu6Yphase by 30% hindering at 300 the°C, movement which was of attributedgrain boundary to the andAl6Cu the6Y deformation phase hindering of matrix the [41movement]. The Al 11ofLa grain3 particles boundary were and very the stable deformation in Al–Cu–La of matrixalloy and [41]. widely The Al distributed11La3 particles in grain were boundaries very stable or dendriticin Al–Cu–La interstices, alloy and which widely effectively distributed hindered in graindislocation boundaries and grain or boundary dendritic migration interstices, during wh highich temperatureeffectively hindered creep [14]. dislocation A large number and of grain small boundaryAl8Cu4Ce migration phases on during dispersoid-free high temperatur grain boundariese creep [14]. have A beenlarge foundnumber in ourof small experimental Al8Cu4Ce alloy phases [1]. onThese dispersoid-free high temperature grain stable boundaries fine particles have canbeen produce found Zenerin our pinning experimental pressure alloy on Zr [1]. dispersoid-free These high temperaturegrain boundaries, stable and fine retard particles its migration can produce during Zene high-temperaturer pinning pressure annealing. on Zr dispersoid-free It is expected that grain the boundaries, and retard its migration during high-temperature annealing. It is expected that the small Al8Cu4Ce dispersed phase mainly distributed at grain boundaries can effectively pin

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Materials 2019, 12, x FOR PEER REVIEW 12 of 14 small Al8Cu4Ce dispersed phase mainly distributed at grain boundaries can effectively pin dislocations and graindislocations boundaries, and grain which boundaries, is beneficial which to the is improvementbeneficial to the of improvement high-temperature of high-temperature tensile deformation uniformitytensile deformation and strength uniformity of Al–Cu–Li and strength alloy. of Al–Cu–Li alloy.

FigureFigure 11. 11.Typical Typical fracture fracture morphologies morphologies of ( a)) Ce-containing Ce-containing and and (b) ( bCe-free) Ce-free alloys alloys after after high high temperature tensile at 100 °C, 250 °C, and 325 °C. temperature tensile at 100 ◦C, 250 ◦C, and 325 ◦C.

4. Conclusions4. Conclusions The inhibitionThe inhibition of theof the main main strength strength phase phase T1T1 coarseningcoarsening through through 0.29 0.29 wt% wt% Ce Ceadditions additions during during high-temperaturehigh-temperature exposure exposure and and the the eff ecteffect of elementof elemen Cet Ce on on the the di ffdiffusionusion of of Cu Cu in in Al–Cu–Li Al–Cu–Li alloyalloy were investigated.were investigated. Moreover, Moreover, the enhancement the enhancement of deformation of deformation uniformity uniformity and and strength strength of of the the Al–Cu–Li Al– Cu–Li alloy due to refined Al8Cu4Ce intermetallic phase in the high-temperature tensile test were alloy due to refined Al Cu Ce intermetallic phase in the high-temperature tensile test were also studied also studied in this8 paper.4 The following conclusions can be drawn. in this paper. The following conclusions can be drawn. (1) The thermal stability of Al–Cu–Li alloy is substantiality improved by Ce addition during (1) Thethermal thermal exposure stability ofat Al–Cu–Lithe medium-high-temperature alloy is substantiality stage improved (170 °C byto Ce270 addition °C). The during enhanced thermal exposurethermal at thestability medium-high-temperature of Ce-containing alloy can stage be (170 attributed◦C to 270 to ◦theC). coarsening The enhanced inhibition thermal of stabilityT1 of Ce-containingthrough impeding alloy the can diffusion be attributed of Cu. Ceto segregates the coarsening at the interface inhibition of the of T1 T1 phase through can restrict impeding the diitsff lateralusion growth, of Cu. Cethereby segregates inhibiting at theits coarsening. interfaceof the T1 phase can restrict its lateral growth, (2)thereby As the inhibiting test temperature its coarsening. increases from room temperature to 325 °C, the elongation of Ce-containing alloy increased rapidly from 7.5 to 17.5%, while that of Ce-free alloy only (2) As the test temperature increases from room temperature to 325 C, the elongation of Ce-containing increased from 7 to 13%. The high tensile strength of Ce-containing◦ alloy is always higher than alloy increased rapidly from 7.5% to 17.5%, while that of Ce-free alloy only increased from 7% to that of Ce-free alloy. The high-temperature stable small Al8Cu4Ce phase in Ce-containing alloy 13%.can The effectively high tensile pin dislocations strength of and Ce-containing grain boundaries, alloy which is always is beneficial higher to than the thatimprovement of Ce-free of alloy. Thethe high-temperature ductility and strength stable of smallthe alloy Al in8Cu the4Ce high-temperature phase in Ce-containing tensile test. alloy can effectively pin dislocations and grain boundaries, which is beneficial to the improvement of the ductility and Authorstrength Contributions: of the alloy Conceptualization, in the high-temperature X.Y.; data curation, tensile J.S.; test. methodology, H.D.; project administration, X.Y.; resources, Z.Z. and D.S. Validation, X.Y.; writing—original draft, X.Y; writing—review and editing, H.D. Authorand Contributions: X.D. Conceptualization, X.Y.; data curation, J.S.; methodology, H.D.; project administration, X.Y.; resources,Funding: This Z.Z. research and D.S. was Validation, funded by (Shandong X.Y.; writing—original Provincial Natural draft, Science X.Y.; Foundation, writing—review China) grant and editing,number H.D. and X.D.(ZR2017PEM005); (A Project of Shandong Province Higher Educational Science and Technology Program) grant Funding:numbersThis (J17KA043, research was J17KB076, funded and by J18KA053). (Shandong Provincial Natural Science Foundation, China) grant number (ZR2017PEM005);Acknowledgments: (A Project We thank of Shandong Caiqiong ProvinceLi, Junfeng Higher Zhao, EducationalWenwen Liu, Scienceand Zhutie and Li, Technology from Yantai Program) Nanshan grant numbers (J17KA043, J17KB076, and J18KA053). University, for contributing to the experimental work in this manuscript. Acknowledgments: We thank Caiqiong Li, Junfeng Zhao, Wenwen Liu, and Zhutie Li, from Yantai Nanshan University, for contributing to the experimental work in this manuscript. Conflicts of Interest: The authors declare no conflict of interest. Materials 2019, 12, 1521 13 of 14

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