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2016-09-07 15:40 [专家论文] 来源于:有色技术与设备信息网
导读:Die cast aluminum alloys are increasingly used in automotive,aerospace and other transportation industries for their light weightand better performance[1–4]. The increasing use ......
Tensile Properties and Deformation Behaviors of a New AluminumAlloy for High Pressure Die Casting
Peng Zhang1, Zhenming Li1,*, Baoliang Liu2, Wenjiang Ding1

1National Engineering Research Center of Light Alloys Net Forming and State Key Laboratory of Metal Matrix Composite, Shanghai Jiao Tong University,Shanghai 200030, China2
National Engineering Research Center of Light Alloys Net Forming, Shanghai 200030, China

1. Introduction
        Die cast aluminum alloys are increasingly used in automotive,aerospace and other transportation industries for their light weightand better performance[1–4]. The increasing use of cast aluminum
components under high stress and high temperature environ-ment has drawn considerable interest in their tensile properties,deformation mechanisms and fracturebehaviors,whicharemainlydetermined by the population of casting defects[5–7]and heat treat-ment conditions[8–12]. For high pressure die casting, the gases haveno enough time to escape from the cavity due to high filling speedand fast cooling rate. As a consequence, these gases in cavity areinevitably involved in the metal liquid, resulting in casting defectssuch as pores forming in the components[13–15]. In addition, turbu-lence flow also leads to the formation of porosity in high pressuredie casting (HPDC). Generally, the castings containing pores cannotundertake heat treatment and also do not work under high tem-perature. The gases in the pores are easy to dilate at the elevatedtemperature environment and cause the surfaces of the castings tobubble, which influences the appearance quality and mechanicalproperties of the products. Heat treatment technologies also causethe changes in size; thus, it is necessary to apply the rectificationwork for the large die cast components with thin wall after heattreatment. Therefore, it would have been more interesting to developnew die cast aluminum alloys that do without heat treatment andprovide higher mechanical properties at room and high tempera-tures so as to widen practical applications of cast aluminum alloys.Current commercial Al–Si based aluminum alloys without suffi-cient ductility (δ < 5%) and better machining performance forstructural applications cannot meet the requirement of modern au-tomobile and aerospace industries[13]. To improve the formability,mechanical strength and ductility of Al–10Si alloys, chemical com-positions of the alloy applied in the present study were designedas follows: (1) Fe content is controlled to as low as 0.15% to avoidforming a large number of brittle intermetallic compounds that de-creasemechanical properties of the castings[16–18]. Mn (0.7%) is usedto replace the Fe and prevent the die-sticking[19]; (2) Besides Cucontent (~1.2%) and Mg content (<0.5%), some other impurity ele-ments (i.e. Zn) are strictly controlled in low level (<0.08%). As anew developed die-casting alloy, the Al–10Si–1.2Cu–0.7Mn alloy (allcomposition in wt% except otherwise stated) 
 

Fig. 1. Diagram of die casting for the standard tensile testing samples of cast alu-
minum alloy according to the specification defined in ASTM B557-06. Sample C cut
from the Al–10Si–1.2Cu–0.7Mn alloy castings was used for tensile testing.
 
Fig. 2. Die casting parameters including filling speed (mm/s) and filling time (s) at
different stages.

is very attractive toautomotive powertrain and structural applications due to its highstrength and ductility as well as corrosion resistance. The tensileperformance of this alloy without experiencing heat treatment canalso meet the requirements for the components. Based on self-strengthening mechanism (aged at room temperature (RT)), the Al–10Si–1.2Cu–0.7Mn alloy exhibits higher hardness (~115 ± 5 (HV-Vickers hardness)), yield strength (~200 ± 10 MPa), ultimate tensilestrength (~330 ± 10 MPa) and ductility (elongation of ~9% ± 1%) (RT).Research on the tensile behaviors of aluminum alloys has indi-cated that their tensile properties, deformation and fracturemechanisms are generally associated with heat treatmentconditions[20–22]and test temperature[23–25]. The present study was
conducted to investigate the effects of natural aging (RT) and testtemperature on the tensile properties, deformation behaviors andfracture mechanisms of the as-cast Al–10Si–1.2Cu–0.7Mn alumi-num alloy so as to widen practical applications of this new die castingalloy.

2. Experimental
2.1. Material and sample preparation
        In this work, high pressure die castings of the Al–10Si–1.2Cu–0.7Mn alloy were used as the study objects. The alloy with a nominalcomposition of Al–10Si–1.2Cu–0.7Mn was prepared from high purityAl, Al–8Si, Al–50Cu, Al–10Mn, and Al–10Ti master alloys in an elec-trical resistance furnace. Before pouring, the melt was held in thefurnace at 690 ± 3 °C for 30 min to ensure homogeneity and disso-lution of the present intermetallic compound. Standard tensilesamples according to ASTM B557-06 (gauge diameter of 6.4 mm andgauge length of 64 mm), as shown in Fig. 1, were obtained by usinga LK400S 4000KN HPDC machine. Die casting parameters includ-ing filling speed (mm/s) and filling time (s) at different stages areshown in Fig. 2. The actual composition was measured by an Optima7300DV inductively coupled plasma optical emission spectrosco-py (ICP-OES) and shown in Table 1. Tensile samples were left on thefloor and at RT for different days (from 0 to 730) before testing theirtensile properties.

2.2. Hardness and tensile testing
        Hardness testing was carried out on a Vickers tester with a loadof 5 kg and holding time of 30 s. No less than 10 indents were mea-sured for each specimen for different natural aging time (from 0.5 hto 17,280 h). Sample C cut from the Al–10Si–1.2Cu–0.7Mn alloy cast-ings was used for tensile testing, as shown in Fig. 1. Tensile testingwas performed on a Zwick/Roell-20kN tensile machine with an at-tached high temperature furnace controlled within ±3 °C. Stress–strain curves were obtained by a knife-edge extensometer (50 mm)attached to the gauge section of the specimens. Test data were usedto calculate the ultimate tensile strength (UTS), 0.2% proof stress(YS), and elongation (A) of the alloy. Different natural aging time(from 0.5 h to 17,280 h, RT) were used to investigate the effects of
the natural aging on the tensile properties of this alloy. All the testswere performed at room temperature (~20 °C) and at a strain rateof ~6.67 × 10−4s−1, and at least four tests were performed for each
natural aging time. Furthermore, five aging time of 2 h (desig-nated as A1), 72 h (A2), 144 h (A3), 1200 h (A4), and 17,280 h (A5) wereselected to obtain the influence of natural aging on the deforma-
tion behavior of the alloy. To study the effects of test temperatureon the tensile properties and deformation behavior of the alloy, seventest temperatures (−70 °C (designated as B1), 20 °C (B2), 150 °C (B
3),200 °C (B4), 250 °C (B5), 300 °C (B6) and 350 °C (B7)) were appliedunder a strain rate of 6.67 × 10−4s−1. The Group-A4(1200 h) sampleswere selected as the object of study because these samples havehigher hardness and tensile properties. High temperature tensilesamples were first heated to the desired temperature and held atthe test temperature for 15 min, and then tested at a strain rate of
6.67 × 10−4s−1. In addition, at least four tensile tests were finishedfor each test temperature.

2.3. Fractographic and microstructural analysis
        The identification of the various phases was carried out by usinga D/MAX-ШA X-ray diffractometer (XRD). Fracture surfaces of the

Table 1 Actual chemical composition of the Al–10Si–1.2Cu–0.7Mn aluminum alloy (wt%)
 

Fig. 3. (a) Microstructure of the tensile bar cross-section for the Al–10Si–1.2Cu–0.7Mn
alloy sample showing band zone. High magnification images of location A (b), lo-cation B (c), and location C (d).
tensile specimens were investigated using a JEOL JSM-6460 fieldemission scanning electron microscope (SEM) with an attachedenergy dispersive spectroscope (EDS) (Japan Electron Optics Ltd.,
Tokyo, Japan). X-ray energy disperse spectrum was used to deter-mine the chemical elements and constituents of the phases. Toinvestigate the microstructures and tensile fracture behaviors of thealloy, metallographic samples were taken from the fracture areasof the tensile samples using a water-cooled low-speed diamond
wheel. Cold mounted samples were carefully ground, polished andetched in a solution composed of 0.5 mL HF and 100 mL H2O. Etchedspecimens for microstructural analysis were observed using a Zeissoptical microscope (equipped with an Axiovision 4.3 Quantimentdigital image analysis system for the quantitative metallography anal-ysis) and a SEM, respectively. The average axis length of α-Al andthe ratio of the major axis length and the minor axis length in opticalmicroscopy (OM) images were measured and calculated using alinear intercept method. Moreover, the area fraction of the α-Al (areaof α-Al/(area of α-Al + area of Si phase particles)) was quantifiedfrom the OM images using an image analysis software.

3. Results and Discussion
3.1. Microstructures
        The typical microstructures of the as-cast Al–10Si–1.2Cu–0.7Mn alloy produced by HPDC are shown in Fig. 3. Low-magnification microstructure observation in Fig. 3(a) clearly indicatesthat the band zone located close to the center can be seen in nearlyall die casting tensile samples. These bands are the commonly-known segregation bands reported in some aluminum andmagnesium die-castingalloys[26–30]. The average distance from theband to the specimen surface and the average width of these bandsare about 1500 m and 200 m, respectively.

        Fig. 3(b–d) shows the typical microstructures in different loca-tions of the as-cast Al–10Si–1.2Cu–0.7Mn alloy tensile bar cross-section (outside of band zone (A), band zone (B) and inside of bandzone (C)). The results indicate that the microstructure of the as-cast Al–10Si–1.2Cu–0.7Mn alloy mainly consists of a soft α-Al matrixand hard Si phase particles (as shown in Fig. 4(b)). A few massiveAlx(Fe/Mn)ySizphase particles (marked by the red arrows in Fig. 4(a),and confirmed by the SEM-EDS analysis in Fig. 4(c) and the resultsof XRD analysis in Fig. 4(d)) can be observed in the as-cast Al–10Si–1.2Cu–0.7Mn alloy specimens. Based on a linear intercept method,it can be seen that the average size of these massive phases is about10–20 m. The average axis length of the α-Al,major axis length(d1) and minor axis length (d2) of the α-Al, ratio of major axis length/minor axis length, and area fraction of the α-Al (≈area of α-Al/(area of α-Al + area of Si phase particles)) in the Al–10Si–1.2Cu–0.7Mn alloy are summarized in Table 2. It is clearly seen that themajor, minor and mean axis length and area fraction of the α-Alin the Al–10Si–1.2Cu–0.7Mn alloy all increase with the locations
changing from outside of the band zone to inside of the band zone.The α-Al on the inside of the band zone exhibits the highest majoraxis length of ~10.7 m, minor axis length of ~7.9 m, mean axislength of ~9.3 m, area fraction of ~68.9%, improved by 43%, 41%,43% and 12%, respectively, in comparison with those of the α-Al onthe outside band. In Table 2, the size of primary alpha-grains in lo-cation C might be underestimated in comparison with that of samegrains viewed in three-dimensional (3D) space. The ratio betweenmajor axis length and minor axis length, however, does not showapparent difference among the different locations of the speci-mens. The average ratio value of α-Al in the Al–10Si–1.2Cu–0.7Mn

Fig. 4. (a) Detailed observation revealing that the alloy specimens contain a few massive phase particles (marked by the red arrows). EDS analysis of location A (b) and
location B (c) in image (a), and XRD curve of the Al–10Si–1.2Cu–0.7Mn alloy (d).

alloy is about 1.34–1.35, indicating that the microstructures of thealloy exhibited apparent spheronization phenomenon during highpressure die casting. As shown in Fig. 3, it can also be seen that the
grain structure of the alloy tensile bar cross-section mainly con-sists of the finer grains and some coarser grains that resulted fromthe externally solidified crystals (ESCs) forming during shotstage
[30–32]. The axis length of α-Al distribution fraction in differ-ent locations of the tensile specimens is shown in Fig. 5. For the Al–10Si–1.2Cu–0.7Mn alloy sample, there is a difference in the ESCsdistribution between the surface and the central region. It is clearlyseen that from surface to center, the density of ESCs increases. Themicrostructures in the central region of cross sections contain a largevolume fraction of coarser and spherical ESCs, and the areas of largerones are more than 700m²(as shown in Fig. 5). In contrast, themicrostructures near the specimen surface contain a significantlylower fraction of ESCs.

        Several studies have been conducted to understand the charac-teristics and the formation mechanism of the segregate bands andthe ESCs in HPDC aluminum and magnesium alloys[26–33]. For hy-poeutectic Al–Si die castings, Laukli et al.[26]pointed out that thecasting conditions and Si content apparently influence the loca-tion and distinctness of the defect bands. Liquid fraction and solutecontent in these bands are higher than those in their surround-ings. Research on the defect bands in HPDC alloys has suggested thatshear deformation of the semi-solid mush[27]and localized defor-mation within partially solid material[28]lead to their formationduring the HPDC processes. In this study, to discern the formationmechanism of these bands, the simulation analysis of moldfilling process using a Flow 3D analysis software was performedand the results were shown in Fig. 6. The result in Fig. 3(a)clearly indicates that the location and symmetry of these bands inthe Al–10Mg–1.2Cu–0.7Mn alloy are fully compatible with thecommonly-known segregation bands reported in some aluminumand magnesium alloys[26–30]. This result is mainly attributed to thesymmetry of the thermal field acting during solidification, whichcan be appreciated by simply looking at the results of the thermalfield analysis, as shown in Fig. 6(a). In a cold-chamber die castingprocess, solidification begins when superheated metal is intro-duced into a relatively cold shot sleeve wall[30]. During the fillingprocess of the casting, the eddy occurs at the bottom of the bar andthe mixture of ESCs and melt tends to fill the cavity in the form ofspiral rise under high pressure and high speed (as shown in Fig. 6(b)).This is due to the specific rheological properties of the filling semi-solid metal[30–33]. Gourlay et al.[30]suggested that the formationmechanism for these segregation bands might be explained basedon the idea of local slip occurring in the dendrite network as a resultof fluid flow past a stagnant solidifying wall layer. The ESC distri-bution can explain that the flow shear induces the migration of ESCstoward the center in the solid and liquid mixture[30–33]
.
Fig. 5. Average axis length of α-Al distribution
fraction of the Al–10Si–1.2Cu–0.7Mn alloy.

Fig. 6. Numerical flow simulation images of thermal field (a) and volume fraction
of entrained air (b) based on Flow-3D analysis software showing the band zone forma-tion process.
3.2. Effects of natural aging time on hardness and tensile properties
        The hardness of the Al–10Si–1.2Cu–0.7Mn alloy is plotted inFig. 7(a), as a function of natural aging time. The result indicatesthat the hardness of the as-cast Al–10Si–1.2Cu–0.7Mn alloy in-creases with the increase of natural aging time. As the aging timeincreases, the hardness value quickly increases up to 116 HV5at 72 h.Further increase of aging time beyond three days, however, doesnot significantly improve the hardness of the alloy. The samples agedat room temperature environment for 17,280 h exhibit the highesthardness of ~121 HV5, an approximate 32% increase compared with
that of the samples left on the floor for 0.5 h (~92 HV5).

        Fig. 7(b) and Table 3 show the RT tensile properties of the Al–10Si–1.2Cu–0.7Mn alloy aged at natural environment (roomtemperature) for different time (from 0.5 h to 17,280 h). The resultsindicate that the YS and UTS of the alloy show trends similar to thatof the hardness with the increase of the natural aging time. In-creasing aging time improves tensile strengths (both YS and UTS)of the alloy,particularly when the aging time is shorter than 72 h.After the natural aging time increases up to 3 days, the further in-crease of time also improves the tensile strengths, but much lesssignificantly. In contrast, it is clearly seen in Fig. 7(b) that the duc-tility of the Al–10Si–1.2Cu–0.7Mn alloy apparently decreases withthe increase of the natural aging time. For the fully as-cast alloy speci-mens kept atroom temperature environment for different time, theelongation varies from 11.6% (0.5 h) to 7.2% (17,280 h). The YS, UTSand elongation of the specimens aged at RT for 72 h areabout 184 MPa, 312 MPa and 10.6%, respectively. In this work, the

Fig. 7. (a) Hardness evolution as a function of aging time (h) during nature aging (RT) and
(b) tensile properties of the Al–10Si–1.2Cu–0.7Mn alloy, as a function of naturalaging time (RT).

Table 3 Tensile properties of the Al–10Si–1.2Cu–0.7Mn alloy with different natural aging time

specimens aged at RT for 17,280 h exhibit the highest tensilestrengths (YS of ~211 MPa and UTS of ~342 MPa) and the lowestductility (elongation of ~7.2%), approximately 27 MPa increase inYS and 30 MPa increase in UTS and ~3.4% decrease in elongation,compared with those of the specimens aged at RT for 72 h (as shownin Table 3). Compared with the as-cast Al–Si–Cu based alloys (such
as A380[13], YS of ~160 MPa, UTS of ~325 MPa, elongation of ~4%),the new alloy shows significant improvements in YS (~51 MPa in-crease), UTS (~17 MPa increase) and elongation (~3.2% increase).
The UTS and ductility of the new alloy are also much higher thanthose of the as-cast Al–Si–Cu–Mg alloys (B390[13]: UTS ≈ 315 MPaand elongation < 1%) and ADC12-0.6Er alloy (UTS ≈ 269 MPa and
elongation ≈ 2%)[34]. Based on self-strengthening mechanism, the newAl–10Si–1.2Cu–0.7Mn alloy is able to obtain higher tensile perfor-mances, which can satisfy the requirements for the automobile andaerospace components. It is generally accepted that the prepara-tion of tensile samples and tensile testing is very costly and time-consuming. Unlike tensile testing, hardness testing is very simple
and quick. This has drawn considerable interest of engineers andscientists to seek some empirical relations between tensile strengthand hardness of metal materials. In the present study, the relation-
ship between hardness (HV5) and ultimate tensile strength (σb) ofthe Al–10Si–1.2Cu–0.7Mn alloy is shown in Fig. 8(a). The result in-dicates that there are good linear correlations between hardness and
ultimate tensile strength. With the best fit of the testing data inFig. 8(a), the linear relationship between HV5and σbof the alloycan be expressed as:
        The comparison between the calculated ultimate tensile strengthbased on Eq. (1) using the hardness (HV5) and the measured actualultimate tensile strength is made in Fig. 8(b). The results show that

Fig. 8. (a) Relationship between the hardness (HV5) and the ultimate tensile strength
(σb) of the Al–10Si–1.2Cu–0.7Mn alloy for different natural aging time and
(b) com-parison between the data based on hardness prediction and the experiment results about the ultimate tensile strength of the alloy.
 

Fig. 9. Tensile properties of the Al–10Si–1.2Cu–0.7Mn alloy
(Group A4– natural agingfor 1200 h), as a function of test temperature.

satisfactory agreement is obtained between the predicted ulti-mate tensile strength and the experimental data with the error bandsof ±5% (R2≈ 0.956). It is therefore believed that based on simple hard-
ness test, we can predict the ultimate tensile strength of this newalloy for high pressure die casting.

3.3. Effects of test temperature on tensile properties
        Fig. 9 and Table 4 show the tensile properties of the Al–10Si–1.2Cu–0.7Mn alloy (Group-A4: 1200 h) tested at different testtemperatures (strain rate of ~6.67 × 10−4s−1). The data of the speci-mens tested at RT (B2) are included in Fig. 9 and Table 4 forcomparison. It is clearly seen that both YS and UTS of the alloy de-crease with the increase of test temperature. In contrast, the ductilityof the alloy significantly increases with increasing test tempera-ture (as shown in Fig. 9 and Table 4). The samples tested at RT (20 °C)exhibit the highest YS (~206 MPa) and UTS (~331 MPa) andthe lowest elongation (~9.8%), improved by 136 MPa in YS and253 MPa in UTS, and reduced by 16.8% in elongation, respectively,compared with those of the samples aged at high temperature (HT)of 350 °C (YS of ~70 MPa, UTS of ~78 MPa, elongation of 26.6%).
 

Table 4Tensile properties of the Al-l OSi-1.2Cu-0.7
Mn alloy tested at different temperatures
 
 
3.4. Strain-hardening behavior and deformation mechanism of thealloy aged at RT for different time
        The effect of natural aging time on the flow behavior of the Al–10Si–1.2Cu–0.7Mn alloy is shown in Fig. 10(a). The result shows thatthe flow stress increases and the total strain (εf) decreases with in-
creasing aging time. It is also clearly seen that the true stress curvesof the alloy aged at RT for different aging time all show hardeningwith the increase of aging time. It is commonly accepted that theworking–hardening behavior of the alloy can be attributed tothe increase of entanglement dislocations and the development of


Fig. 10. (a) Effect of natural aging time on true stress–logarithmic strain curves of the Al–10Si–1.2Cu–0.7Mn alloy (A1– 2 h, A2– 72 h, A3– 144 h, A4– 1200 h and A5–17,280 h). (b) The magnified curve of location B in (a) showing the Portevin–Le Chatelier effect. σ1is stress drop amplitude, σ2is stress increase amplitude; t1is stress de-crease time, t2is reloading time. (c) True stress–logarithmic plastic strain curves showing the effect of natural aging time on flow behavior of the alloy. (d) Strain-hardeningrate measured at a plastic strain of 0.0015, as function of natural aging time for the alloy. (e) Determination of n and k values for the alloys tested at different natural agingtime by linear fit to the log (true stress)–log (logarithmic plastic strain) curves. The true plastic strain ranges from about 0.01 up to instability. (f) n and (g) k measured ata logarithmic plastic strain range from about 0.01 up to instability, as function of natural aging time for the alloy.

the dislocation multiplication. The result in Fig. 10(b) clearly indi-cates that the true stress–logarithmic strain curves of the samplesaged at RT for 2 h and tested at RT exhibit the Portevin–Le Chatelier
(PLC) effect confirmed as Type A behavior[35–41]. The PLC effect as oneof the most prominent examples of instability often occurs in somealuminum alloys under certain range of test temperature and strainrate. Discontinuous yielding in the curves is considered as a conse-quence of dynamic strain aging (DSA)[42–48]. In this work, the PLCbehavior for the alloy aged at RT for a shorter time (<36 h) is prob-ably attributed to the interactions between the Si and Cu atoms insolution and the glide dislocations. With increasing the natural agingtime (>36 h), the PLC effect disappears because it is difficult to formthe effective solution atmospheres locking mobile dislocations. It isalso seen that the curves in Fig. 10(a) mainly consist of two regimesincluding linear hardening stage and parabolic hardening stage. Thetransition between two regimes is likely controlled by the relativestrength of the soft matrix and the harder phase particles[49–54].

        Generally, the stress–strain curves of aluminum alloys in plasticdeformation stage can be expressed by using the Hollomon equa-tion (σ = kεn). The true stress–logarithmic plastic strain curves of theas-cast Al–10Si–1.2Cu–0.7Mn alloy aged at RT for different time (from0.5 h to 17,280 h) and tested at RT are shown in Fig. 10(c). The effectof natural aging time on the strain-hardening rate (dσ/dε) at a lowstrain of 0.0015 is shown in Fig. 10(d). It can be seen that the dσ/dε(from 12 to 8.8 GPa) at low strains decreases with the improve-ment of aging time. To obtain the strain-hardening exponent (n) and
the strain-hardening coefficient (k) of the alloy aged and tested atRT, the log (true stress)–log(logarithmic plastic strain) data over astrain range from about 0.01 up to instability are plotted in Fig. 10(e).
By fitting the linear relationship, the constant values of n and k aredetermined. Fig. 10(f–g) shows the influences of natural aging timeon the n value and the k value at the plastic strain from 0.01 to in-
stability, respectively. The results indicate that with the increase ofaging time, n value decreases (from 0.203 to 0.155) and k value in-creases (from 468 to 533). The alloy aged at RT for two years and
tested at RT exhibits the lowest n value (~0.155) and the highest k(~533 MPa). For the as-cast Al–10Si–1.2Cu–0.7Mn alloy aged at RTfor 2 h, the Al matrix is relatively soft and the interaction amongdislocations, eutectic phase particles, solid solution atoms and grainboundaries dominates the hardening behavior, tensile propertiesand fracture mechanism. The results in Fig. 7 clearly indicate thatwith increasing the natural aging time, the hardness and the tensilestrength (both YS and UTS) of the Al–10Si–1.2Cu–0.7Mn alloy onlyexhibit the first peak values. In this case, it is hard to observe obviousprecipitation spots in electron diffraction pattern[55–57]. This is mainlybecause at the lower natural aging temperature (RT), intracrystallineprecipitation is inadequate[55–57]. It is believed that the strengthen-ing effect on the alloy mainly comes from the solute atoms, Guinier–Preston (GP) zone, pure Al, grain boundary as well as phase particle.Therefore, for the Al–10Si–1.2Cu–0.7Mn alloys, increasing naturalaging time tends to lead to decreases in n value and ductility andimprovements in k value and tensile strengths (both YS and UTS).

3.5. Strain-hardening behavior and deformation mechanism of thealloy tested at different temperatures
        The effect of test temperature on the flow behavior of the Al–10Si–1.2Cu–0.7Mn alloy is shown in Fig. 11(a). The curve of thesamples tested at RT is also included in the figure for comparison.It is clearly seen that the flow stress of the alloy decreases and thetotal strain increases with increasing test temperature. The resultsin Fig. 11(a) also indicate that the true stress–logarithmic strain curves
all show softening after the flow stress reaches a maximum withincreased logarithmic strain. In addition, as shown in Fig. 11(b), itcan also be seen that the PLC behavior only occurs in the Al–10Si-1.2Cu–0.7Mn alloy samples tested at HT of 250 °C. According to thestress serrations or band propagation characteristics, the PLC effectin this work is considered as Type C behavior[35–41]. Based on theinteraction between dislocations and obstacles (such as precipi-tates), dynamic strain aging mechanism can be used to explain thePCL behavior in the samples tested at 250 °C[51]. In the present study,the mechanism of PLC effect in the samples tested at RT and HT of250 °C is quite different. For the alloy tested at RT, the interactionsbetween solute atoms and dislocations determine PLC effect. In con-trast, when the test temperature is 250 °C, the precipitated phasessignificantly influence PLC behavior in the alloy. The true stress–logarithmic plastic strain curves of the Al–10Si–1.2Cu–0.7Mn alloytested at different test temperatures (from −70 °C to 350 °C) and thesame strain rate (~1.67 × 10−4s−1) are shown in Fig. 11(c). The effectof test temperature on the strain-hardening rate (dσ/dε) at a lowstrain of 0.0015 is shown in Fig. 11(d). The result shows that for thealloy tested at different temperatures, the improvement of test tem-perature significantly reduces the strain-hardening rate dσ/dε at lowstrains (from 14.2 to 2.26 GPa), corresponding to the decreased yieldstrengths. The n and k values were also determined by a linear fitto the log(true stress)–log(logarithmic plastic strain) data over a strainrange from about 0.01 up to instability, as shown in Fig. 11(e). Theinfluences of test temperature on the n and k values of the Al–10Si–1.2Cu–0.7Mn alloy are shown in Fig. 11(f–g). It is clearly seen
that increasing test temperature (from −70 °C to 350 °C) signifi-cantly reduces n (from 0.172 to 0.027) (Fig. 11(f)). The results inFig. 11(g) indicate that the k of the alloy also decreases with in-creasing test temperature. For example, the specimens tested at RTexhibit the highest k value of ~513 MPa, improved by ~426 MPa, com-pared with that of the samples tested at 350 °C (87 MPa). For theAl–10Si–1.2Cu–0.7Mn alloy tested at HT, besides the matrix strength-ening (contribution from pure Al, solid solution atom, precipitationand grain boundary), only eutectic phase particles contribute to tensileproperties and working-hardening. Therefore, the differences betweentensile properties and strain-hardening behavior of the alloy testedat low temperature (LT of −70 °C), RT and HT can be greatly attrib-uted to different matrix strengths and phase particle strengths ofthe alloy under different temperature environments. During tensiletest, accumulation and annihilation of dislocations are consideredas the possible mechanisms leading to different work-hardening be-haviors and tensile properties of the alloy[58]. Increasing testtemperature can accelerate the annihilation processes correspond-ing to lower n and k values and tensile strengths. In addition, it canbe seen from Fig. 4(a) that the Al–10Si–1.2Cu–0.7Mn alloy containssome massive Alx(Fe/Mn)ySizcompounds. Generally, the big par-ticles acting as barriers for dislocation motion tend to constrain thedeformation of the matrix alloy during tensile test[58]. These com-pounds, however, are easy to soften at elevated temperature(>130 °C)[59,60], leading to the decreases of the HT tensile strengths(both YS and UTS), n and k values and the increase of the HT duc-tility of the alloys. Thus, compared with the samples tested at LTand RT, the Al–10Si–1.2Cu–0.7Mn alloy samples tested at HT showapparently higher tensile strengths, strain-hardening exponent andcoefficient, and lower ductility.

3.6. Fractographic observations
        The typical fractographs of the Al–10Si–1.2Cu–0.7Mn alloy testedat RT (20 °C), 150 °C, 250 °C and 350 °C are shown in Fig. 12. Theresult in Fig. 12(a) clearly indicates that for the samples tested at
LT and RT, the fracture surfaces are mainly composed of tear ridgesas well as cracked and debonded Al–Si eutectic particles. It can alsobe seen that there are some dimples on the fracture surface of thealloy samples (as shown in Fig. 12(b)), corresponding to the highductility of these samples (elongation > 9.5%). It is interesting to note

Fig. 11. (a) Effect of test temperature on true stress–logarithmic strain curves of the Al–10Si–1.2Cu–0.7Mn alloy (B1: −70 °C, B2: 20 °C, B3: 150 °C, B4: 200 °C, B5: 250 °C, B6:300 °C and B7: 350 °C). (b) The magnified curve of location B in (a) showing the Portevin–Le Chatelier effect. (c) True stress–logarithmic plastic strain curves showing theeffect of test temperature on flow behavior of the alloy. (d) Strain-hardening rate measured at a plastic strain of 0.0015, as function of test temperatures for the alloy. (e)Determination of n and k values for the alloys tested at different test temperatures by linear fit to the log (true stress)–log (logarithmic plastic strain) curves. The logarith-mic plastic strain ranges from about 0.01 up to instability. (f) n and (g) k measured at a logarithmic plastic strain range from about 0.01 up to instability, as function of testtemperatures for the alloys.
 
 
that the size of these dimples is similar to that of the Alx(Fe/Mn)ySizphase particles, indicating that cracking and debonding ofthese particles contribute to the crack initiation, crack propaga-tion and final fracture during tensile test. In addition, it is also clearlyseen from Fig. 12(c–f) that the fracture surface of the tensile samplestested at HT shows a feature similar to that of the samples testedat LT and RT. Compared with the fracture surface of the samplestested at LT and RT, more deep dimples are observed on that of thesamples tested at HT, corresponding to significant improvementsin ductility at elevated temperature environment. Therefore, it isrational to suggest that cracking and debonding of phase particles(both Si and Alx(Fe/Mn)ySiz) are the main damage mechanisms priorto the final fracture of the Al–10Si–1.2Cu–0.7Mn alloy tested at LT,RT and HT.
 
        Generally, high speed should cause gas entrapment during moldfilling (as shown in Fig. 6) and porosities in these die castings. Inthis work, however, no porosities and inclusions were observed on
the fracture surface of the tensile specimens, as shown in Fig. 13.This is due to the relatively perfect designs of the slag ladle and theexhaust slot in the present study. Therefore, it is believed that the
tensile properties and deformation behavior of the Al–10Si–1.2Cu–0.7Mn alloy are not influenced by casting defects.

3.7. Tensile instability and fracture mechanism
        Generally, the global tensile instability of a material can be de-termined by using Considère criterion[54], where the strain-hardeningrate (dσ/dε) is equal to the flow stress (σ). In other words, the n of
Fig. 12. SEM fractographs showing the effect of test temperature on the fracture in the Al–10Si–1.2Cu–0.7Mn alloy: (a, b) 20 °C; (c, d) 150 °C; (e, f) 300 °C. The Alx(Fe, Mn)ySizphase particles and Sieutectic particles are marked by red and blue arrows, respectively.

the alloy is also the same as the corresponding critical strain (εi)when global instability takes place. The tensile instability behav-iors of the Al–10Si–1.2Cu–0.7Mn alloy aged at RT for different timeand tested at RT are shown in Fig. 14(a). It is clearly seen that forthe alloy aged and tested at RT, the flow stress curves do not in-tersect with the strain-hardening curves prior to failure. Fig. 14(b)shows the tensile instability behaviors of the Al–10Si–1.2Cu–0.7Mn alloy tested at different temperatures. The result indicatesthat for the alloy tested at HT except the samples tested at 150 °C,the strain-hardening curves do in fact intersect the flow stress curvesprior to failure. The results in Fig. 15(a) indicate that the ratios oftensile instability (εi= n) to fracture strain (εf) for the alloy testedat RT are close to or larger than 1. This result suggests that for theAl–10Si–1.2Cu–0.7Mn alloy, tensile failure occurs before the globalinstability takes place. Premature damage and high damage rate (par-ticle cracking) occurring in the alloy tested at RT are mainlyattributed to localized deformation of the inhomogeneous micro-structure. Moreover, it is also seen from Fig. 15(a) that increasingnatural aging time slightly improves the ε
i/εfratio of the alloy (from2.055 to 2.402). As shown in Fig. 15(b), it can also be seen that theεi/εfratios of the alloy tested at HT are lower than 1, indicating thatthe global instability occurs at strains apparently below the true frac-ture strains (εi> n). Compared with LT and RT, high temperatureenvironment softens the matrix and eutectic phase particles andreduces the damage accumulation rate in the alloy. These resultsindicate that there is still a post necking damage[54]. Furthermore,it is also clearly seen that the ratio εi/εfof the alloy decreases withthe increase of test temperature (from 2.795 to 0.125) (Fig. 15(b)).Like A356/A357 aluminum castings[54], both local and global insta-bilities determine the fracture behavior of the Al–10Si–1.2Cu–0.7Mn alloy. The macroscopic fracture morphologies of the alloyspecimens tested at different temperatures are shown in Fig. 16. Itis clearly seen that shearing occurs in the alloy tested at LT and RT,corresponding to lower ductility with a lot of tear ridges and deepdimples. On the other hand, necking is observed in the alloy testedat HT and the failure features are macroscopic shear fracture witha large number of shallower shear dimples and some big and deepdimples on the microscale. In the alloy tested at RT, the critical

Fig. 13. Lower magnification images of (a) Fig. 12(a), (b) Fig. 12(b) and (c) Fig. 12(c).
Fig. 14. Tensile instability plots for the Al–10Mg–1.2Cu–0.7Mn alloy for different natural aging time and tested at different test temperatures: (a) A1– 2 h, A2– 72 h, A3–144 h, A4– 1200 h and A5– 17,280 h;
(b) B1: −70 °C, B3: 150 °C, B4: 200 °C, B5: 250 °C, B6: 300 °C and B7: 350 °C.
Fig. 15. Ratio of tensile instability strain εito fracture strain εfof the Al–10Si–1.2Cu–0.7Mn alloy for different natural aging time (a) and tested at different test tempera-tures (b). εi/εf> 1 indicates that local fracture governs tensile instability, while εi/εf< 1 indicates more uniform damage accumulation where tensile instability occurs at theonset macroscopic necking.

amount of damage for failure is easily reached prior to global necking,which is attributed to the large and hard particles. In contrast, it isdifficult to obtain the critical amount of damage before the global
necking occurs in the alloy tested at HT. This result is mainly at-tributed to the soft matrix and particles as well as much lowerdamage rate during high temperature test.

4. Conclusions
        This paper investigates the influences of natural aging time andtest temperature on the tensile properties, deformation behaviorand fracture mechanism of a recently developed high-ductility castaluminum alloy Al–10Si–1.2Cu–0.7Mn for high pressure die casting.Based on the results obtained in this work, the following conclu-sions can be drawn.
 
        (1) Based on self-strengthening mechanism, the Al–10Si–1.2Cu–0.7Mn alloy aged and tested at RT can provide high hardness(HV5: 92–122), yield strength (162–212 MPa), ultimate tensilestrength (250–342 MPa), and ductility (elongation of 7.52%–11.56%), much higher than those of some current commercialaluminum alloys such as A380.
Fig. 16. Macroscpic tensile fracture features of the Al–10Si–1.2Cu–0.7Mn alloy testedat different temperatures (B2: 20 °C, B3: 150 °C, B4: 200 °C, B5: 250 °C, B6: 300 °C andB7: 350 °C).

        (2) The hardness, YS and UTS increase and the ductility of theAl–10Si–1.2Cu–0.7Mn alloy decreases with the increase of thenatural aging time. After the aging time increases up to 72 h,the further increase of time also improves the tensile strengths,but much less significantly. With the increase of test tem-perature (from −70 °C to 350 °C), the Al–10Si–1.2Cu–0.7Mnalloy shows apparent decreases in YS (162–70 MPa) and UTS(342–78 MPa) and improvements in ductility (elongation of7.52%–26.6%). Hardness can be used to predicate the ulti-mate tensile strength of this new alloy.

        (3) For the alloy tested at RT, the Portevin–Le Chatelier effect takesplace due to the interactions between solid solution atomsand dislocations. In contrast, the PLC behavior in the alloytested at HT of 250 °C can be explained by using dynamicstrain aging mechanism. With increasing the natural agingtime, the n value increases and the k value decreases. The nand k values of the alloy apparently decrease with the in-crease of test temperature.

        (4) The tensile failure of the Al–10Si–1.2Cu–0.7Mn alloy is origi-nated from the cracked and debonded eutectic phase particles.Both local instability (RT, shearing) and global instability (HT,necking) determine the fracture behavior of the Al–10Si–1.2Cu–0.7Mn alloy. The tensile properties, deformationbehavior and fracture mechanism of the alloy tested at RT (LT)and HT mainly depend on the matrix strengths, eutectic phaseparticle strengths and damage rate.

Acknowledgments
        This work was supported by the Project Funded by China Post-doctoral Science Foundation (No. 2015M571562). The authors aregrateful to Prof. Liming Peng (Shanghai Jiaotong University) for hishelpful discussion.
 

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