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Please note: Images will appear in color online but will be printed in black and white.ArticleTitle Improving High-Temperature Tensile and Low-Cycle Fatigue Behavior of Al-Si-Cu-Mg Alloys Through

Micro-additions of Ti, V, and ZrArticle Sub-TitleArticle CopyRight The Minerals, Metals & Materials Society and ASM International

(This will be the copyright line in the final PDF)Journal Name Metallurgical and Materials Transactions ACorresponding Author Family Name Czerwinski

ParticleGiven Name FrankSuffixDivisionOrganization CanmetMATERIALS/CanmetMATRIAUX, Natural Resourses CanadaAddress 183 Longwood Road South, Room 259C, L8P 0A5, Hamilton, ON, CanadaEmail [email protected]

Author Family Name ShahaParticleGiven Name S. K.SuffixDivision Department of Mechanical and Industrial EngineeringOrganization Ryerson UniversityAddress 350 Victoria Street, M5B 2K3, Toronto, ON, CanadaEmail

Author Family Name KasprzakParticleGiven Name W.SuffixDivisionOrganization CanmetMATERIALS/CanmetMATRIAUX, Natural Resourses CanadaAddress 183 Longwood Road South, Room 259C, L8P 0A5, Hamilton, ON, CanadaEmail

Author Family Name FriedmanParticleGiven Name J.SuffixDivision Department of Mechanical and Industrial EngineeringOrganization Ryerson UniversityAddress 350 Victoria Street, M5B 2K3, Toronto, ON, CanadaEmail

Author Family Name Chen

ParticleGiven Name D. L.SuffixDivision Department of Mechanical and Industrial EngineeringOrganization Ryerson UniversityAddress 350 Victoria Street, M5B 2K3, Toronto, ON, CanadaEmail

ScheduleReceivedRevisedAccepted

Abstract High-temperature tensile and low-cycle fatigue tests were performed to assess the influence of micro-additions of Ti, V, and Zr on the improvement of the Al-7Si-1Cu-0.5Mg (wt pct) alloy in the as-castcondition. Addition of transition metals led to modification of microstructure where in addition toconventional phases present in the Al-7Si-1Cu-0.5Mg base, new thermally stable micro-sized Zr-Ti-V-richphases Al21.4Si4.1Ti3.5VZr3.9, Al6.7Si1.2TiZr1.8, Al2.8Si3.8V1.6Zr, and Al5.1Si35.4Ti1.6Zr5.7Fe were formed. Thetensile tests showed that with the increase in the testing temperature from 298 K to 673 K (25 C to400 C), the yield stress and tensile strength of the present studied alloy decreased from 161 to 84 MPa andfrom 261 to 102 MPa, respectively. Also, the studied alloy obtained 18, 12, and 5 pct higher tensilestrength than the alloy A356, 354 and existing Al-Si-Cu-Mg alloy modified with additions of Zr, Ti, andNi, respectively. The fatigue life of the studied alloy was substantially longer than those of the referencealloys A356 and the same Al-7Si-1Cu-0.5Mg base with minor additions of V, Zr, and Ti in the T6condition. Fractographic analysis after tensile tests revealed that at the lower temperature up to 473 K(200 C), the cleavage-type brittle fracture for the precipitates and ductile fracture for the matrix weredominant while at higher temperature fully ductile-type fracture with debonding and pull-out of crackedparticles was identified. It is believed that the intermetallic precipitates containing Zr, Ti, and V improvethe alloy performance at increased temperatures.

Footnote Information Published with permission of Her Majesty the Queen in Right of Canada pertains to F. Czerwinski andW. Kasprzak.Manuscript submitted December 22, 2014.

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1

2 Improving High-Temperature Tensile and Low-Cycle Fatigue

3 Behavior of Al-Si-Cu-Mg Alloys Through Micro-additions

4 of Ti, V, and Zr5

6 S.K. SHAHA, FRANK CZERWINSKI, W. KASPRZAK, J. FRIEDMAN, and D.L. CHEN78 High-temperature tensile and low-cycle fatigue tests were performed to assess the influence of9 micro-additions of Ti, V, and Zr on the improvement of the Al-7Si-1Cu-0.5Mg (wt pct) alloy in10 the as-cast condition. Addition of transition metals led to modification of microstructure where11 in addition to conventional phases present in the Al-7Si-1Cu-0.5Mg base, new thermally stable12 micro-sized Zr-Ti-V-rich phases Al21.4Si4.1Ti3.5VZr3.9, Al6.7Si1.2TiZr1.8, Al2.8Si3.8V1.6Zr, and13 Al5.1Si35.4Ti1.6Zr5.7Fe were formed. The tensile tests showed that with the increase in the testing14 temperature from 298 K to 673 K (25 C to 400 C), the yield stress and tensile strength of the15 present studied alloy decreased from 161 to 84 MPa and from 261 to 102 MPa, respectively.16 Also, the studied alloy obtained 18, 12, and 5 pct higher tensile strength than the alloy A356,17 354 and existing Al-Si-Cu-Mg alloy modified with additions of Zr, Ti, and Ni, respectively. The18 fatigue life of the studied alloy was substantially longer than those of the reference alloys A35619 and the same Al-7Si-1Cu-0.5Mg base with minor additions of V, Zr, and Ti in the T6 condition.20 Fractographic analysis after tensile tests revealed that at the lower temperature up to 473 K21 (200 C), the cleavage-type brittle fracture for the precipitates and ductile fracture for the matrix22 were dominant while at higher temperature fully ductile-type fracture with debonding and pull-23 out of cracked particles was identified. It is believed that the intermetallic precipitates containing24 Zr, Ti, and V improve the alloy performance at increased temperatures.

2526 DOI: 10.1007/s11661-015-2880-x27 The Minerals, Metals & Materials Society and ASM International 201528

29 I. INTRODUCTION

30 TO develop fuel-efficient vehicles, weight reduction31 using light-weight materials is one of the several32 strategies pursued.[13] Al and its alloys, especially Al-33 Si-Cu-Mg grades, are light-weight metallic materials34 commonly used in the automotive and aerospace indus-35 tries. The major factor limiting automotive applications36 of Al alloys is their thermal stability. At temperatures37 exceeding 473 K (200 C), phases such as Al2Cu, Mg2Si,38 and/or Al2CuMg, which maintain the alloy strength,39 usually coarsen or dissolve. This results in reduced40 performance, consequently limiting their practical41 applications in engine blocks, cylinder heads, or heat42 shields.[47] To meet the stringent industrial require-43 ments, properties of Al-Si-Cu-Mg alloys should remain44 stable up to at least 573 K (300 C).[8]

45 Among many ways of improving the high-tem-46 perature properties, alloying with transition metals to

47form thermally stable and coarsening-resistant pre-48cipitates was found to be very promising. To achieve49this, different alloying elements such as Ni, Fe, Cr, Ti, V,50and Zr in cast Al-Si alloys were tested in the lit-51erature.[7,926] The influence of Zr and Ti additions in Al52alloys was a subject of a number of studies with major53conclusions that the morphology and type of phases54formed during either casting or heat treatment control55the high-temperature properties.[16,19] Mahmudi et al.[11]

56and Sepehrband et al.[10] modified the A319 Al-Si cast57alloy with the addition of Zr and improved its tensile58strength and wear resistance. Research also showed that59V could enhance the alloy performance by forming Al3V60or Al10V.

[24] Recently, Mohamed et al. and others[7,27]

61showed that addition of Ni in the Al-Si-Cu-Mg alloy62caused a reduction in alloy strength at room tem-63perature mainly due to a decrease in the available Cu for64precipitation strengthening through forming Al3CuNi.65In this alloy, an increase in content of the Al3CuNi and66Al9NiFe phases was responsible for some reduction in67ductility. At the same time, the presence of Fe in the Al-68Si-Cu-Mg alloy led to the formation of the Fe-contain-69ing b-Al5FeSi phase, which is also responsible for the70reduction of alloy ductility. Our previous studies[14,2830]

71showed that addition of Ti-V-Zr in Al-Si-Cu-Mg alloys72did not give rise to any copper-containing phases with73Ti-V-Zr but rather modified the Fe-containing b-Al5Fe-74Si phases, thereby improving the tensile/compression75strength and low cycle fatigue (LCF) strength in the T676heat-treated condition. However, there are no studies on

S. K. SHAHA, J. FRIEDMAN, and D.L. CHEN, are with theDepartment of Mechanical and Industrial Engineering, RyersonUniversity, 350 Victoria Street, Toronto, ON M5B 2K3, Canada.FRANK CZERWINSKI, P. Eng., Group Leader, Senior ResearchScientist and W. KASPRZAK, are with the CanmetMATERIALS/CanmetMATERIAUX, Natural Resourses Canada, 183 LongwoodRoad South, Room 259C, Hamilton, ON L8P 0A5, Canada. Contacte-mail: [email protected]

Published with permission of Her Majesty the Queen in Right ofCanada pertains to F. Czerwinski and W. Kasprzak.

Manuscript submitted December 22, 2014.

METALLURGICAL AND MATERIALS TRANSACTIONS A xxx1

Journal : MMTA Dispatch : 24-3-2015 Pages : 16

PIPS No. : 2880h LE h TYPESET

MS Code : h CP h DISK

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77 tensile properties at higher temperatures and LCF78 fatigue resistance of Al-Si-Cu-Mg as-cast alloys mod-79 ified with addition of Zr, Ti, and V. Thus, it is not clear80 how the Zr-Ti-V-rich phases present in the as-cast state81 behave during LCF at room temperature and especially82 during deformation at increased temperatures.83 Therefore, the aim of this study was to investigate84 deformation mechanisms of the Al-Si-Cu-Mg alloy with85 addition of Zr, Ti, and V at various temperatures under86 tensile loading and under LCF with different strain87 amplitudes. The results are of importance since die88 castings are considered as finished products and often89 they are not subjected to subsequent heat treatment.

90 II. EXPERIMENTAL

91 A. Alloy Processing and Microstructure

92 The studied alloy was prepared by melting several93 master alloys of Al-Cu, Al-Si, Al-Ti, Al-Zr, and Al-V94 along with pure Mg and Al. The targeted alloy95 chemistry is given in Table I. First, pure Al was melted96 under a protective atmosphere. During addition of97 master alloys, the melt temperature was increased to98 1073 K to 1123 K (800 C to 850 C) to ensure their99 proper dissolution. The alloy was cast at 1013 K100 (740 C) into the wedge-shaped steel mold preheated101 to a temperature of 673 K (400 C). Further casting102 details are explained in Reference 29. After casting,103 samples were collected from the middle of the wedge104 having SDAS of 25 lm. The alloy samples for metal-105 lographic investigations were prepared following a106 standard metallographic technique.[29] The microstruc-107 ture was examined in an unetched condition using an108 optical microscope (OM) equipped with a quantitative109 image analyzer (CLEMEX software), and a scanning110 electron microscope (SEM) coupled with energy-disper-111 sive X-ray spectroscopy (EDX). The deep etching112 was performed using 5 pct NaOH as described in113 Reference 30.

114 B. Measurement of Tensile and Fatigue Properties

115 The extraction of samples from the wedge having116 secondary dendrite arm spacing (SDAS) of 25 lm was117 illustrated in our previous study.[31] The tensile testing118 was performed on sub-size rectangular bar samples with119 a gage length of 25 mm (or parallel length of 32 mm)120 and a cross section of 6 9 6 mm2 following the ASTM:121 E8M-11 standard at temperatures of 298 K (25 C)122 (room temperature), 473 K, 573 K, and 673 K (200 C,123 300 C, and 400 C) with a deviation of 5 C at a124 strain rate of 103 s1 using a computerized United125 Tensile Testing machine (Model: STM, 50 kN). The126 strain was measured by a clip-on 25 mm extensometer

127attached to the gage length. Prior to testing, each128specimen was kept in the heating chamber of the testing129machine for 10 minutes at the desired temperature. The130temperature of the heating chamber was precisely131measured in two ways; using K-type thermocouple132which was placed near the sample gripe with multi-meter133and pre-assemble the heating chamber thermocouple.134After testing, the specimens were quenched in water to135keep the same microstructure at testing temperature. As136described in Reference 32 after tensile test, the calculat-137ed true stress versus true strain graphs were plotted138using collected raw data. From the graph, 0.2 pct offset139was calculated and considered as yield stress/strength140(YS) of the materials, while maximum stress was141considered as ultimate tensile strength (UTS). The YS142and UTS values of average of at least two samples were143plotted with respect to the testing temperature.144For cyclic testing, a similar type of sub-size bar145sample was used. The samples were first polished by146SiC sand paper up to 600 grade to remove the147machining effect. The strain-controlled, pull-push-type148fatigue tests were conducted in accordance with the149ASTM: E606 standard at room temperature with a15025 mm extensometer using a computerized Instron1518801 fatigue testing machine operated by Bulehill152LCF2 software. Triangular waveform loading with a153zero mean strain (i.e., a strain ratio of Re = 1,154completely reversed strain cycle) at a constant strain155rate of 1 9 102 s1 was applied during cyclic defor-156mation tests. The cyclic frequency was varied depend-157ing on the strain amplitude to maintain a fixed strain158rate. The strain was also measured by a clip-on 25 mm159extensometer attached to the gage length. Low-cycle160fatigue tests were performed at total strain amplitudes161of 0.1, 0.2, 0.3, 0.4, 0.5, and 0.6 pct with at least two162samples tested at each level of strain amplitude. If the163sample survived 10,000 cycles, then the strain-con-164trolled tests were transferred to load control at a165frequency of 50 Hz using a sinusoidal waveform. The166fatigue life was considered as the number of cycles to167completely separate apart of the samples. The fracture168surfaces of the tensile and fatigue specimens were169examined via scanning electron microscope (SEM) to170identify fatigue crack initiation sites and propagation171characteristics.

172III. RESULTS

173A. Microstructure

174The Al-Si-Cu-Mg alloy established a complex mi-175crostructure as reported in many studies. The presence176of the alloying elements Zr-Ti-V in this study led to the177development of an even more complex microstructure.178Typical OM and SEM microstructures of the studied

Table I. Chemical Composition of Al-Si-Cu-Mg Alloy Modified with Addition of Ti, V, and Zr in wt pct

Si Cu Mg Fe Sr Mn Zr Ti V Al

7.02 0.95 0.48 0.090 0.012 0.005 0.47 0.20 0.32 bal.

2xxx METALLURGICAL AND MATERIALS TRANSACTIONS A




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179 alloy in the unetched and deep-etched conditions are180 shown in Figure 1. The chemical composition of phases181 identified using EDX point analysis along with the182 literature suggestions are given in Table II. The pre-183 sented chemistries are the average values of at least five184 measurements for each phase. It is seen that Zr always185 forms intermetallics with other transition metals (Ti and186 V) and increasing the Zr content in the Zr-Ti-V-rich187 phases leads to reduced content of Si and V. The188 metallographic analysis revealed that the alloy189 microstructure consisted of a-Al dendrites (#1),190 fibrous-like modified Al-Si eutectic (#2) and eight (8)191 distinct intermetallic phases as listed in Table II and192 Figures 1(a) through (c). As calculated from the phase193 chemistry, two types of Cu-rich phases are Al2.1Cu (#3)194 and Al8.5Si2.4Cu (#4), which are generally suggested as195 h-Al2Cu

[7] and Al-Al2Cu-Si ternary eutectic[33] phase,

196 respectively. At the same time, the Mg- and Fe-rich197 phases are calculated as Al7.2Si8.3Cu2Mg6.9 (#5) and198 Al14Si7.1FeMg3.3 (#6) where those phases are suggested199 as Q-Al5Cu2Mg8Si6 and p-Al8FeMg3Si6 phases, respec-200 tively. Similar types of phases were found in commercial201 Al-Si-Cu-Mg alloys in the as-cast condition.[7,27,34,35]

202 The Zr-Ti-V-containing phases which are unique for203 this study are also identified in the alloy microstructure.204 There are two morphologies: plate and bulk shape.205 Those phases are frequently noted in the alloy mi-206 crostructure as seen in Figure 1(c). The calculated Zr-Ti-207 V-containing phases are Al21.4Si4.1Ti3.5VZr3.9 (#7),208 Al6.7Si1.2TiZr1.8 (#8), and Al2.8Si3.8V1.6Zr (#9) which209 were designated in the literature as (AlSi)3(TiVZr),210 (AlSi)3(TiZr), and (AlSi)2(VZr), respectively.

[7,14,27,29]

211 As reported in Reference 30, the Zr-Ti-V-containing212 phases are thermally stable which is beneficial to213 improve the alloy high-temperature performance. The214 identified intermetallic phases were found most often in215 interdendritic regions. As observed in Figure 1(c), the216 plate-shaped Zr-Ti-V-containing phases grew not only217 in interdendritic regions but also across a-Al grains and218 nucleated during alloy solidification. Another valuable219 modification of alloy microstructure was observed for220 the Fe-containing b-Al5FeSi phase which is generally221 needle-like in shape and detrimental for the alloy222 strength and ductility. However, it is assumed that the223 Fe-containing b-Al5FeSi phase turns into the Al5.1Si35.4-224 Ti1.6Zr5.7Fe (#10) phase which could be described as225 (AlSi)2(TiZr)Fe, as reported in References 13, 14.226 Similar types of intermetallics, its formation mechanism,227 and phase morphology were reported in detail in our228 previous studies.[13,14,29,36]

229 B. Mechanical Properties

230 1. Alloy tensile properties231 The yield strength (YS) and ultimate tensile strength232 (UTS) of the modified alloy, obtained at different233 temperatures at a strain rate of 103 s1, are plotted in234 Figure 2 and compared with Al-Si alloys 354, A356 and235 the Al-Si-Cu-Mg alloy modified with the addition of Ti,236 Zr, and Ni. As seen in Figure 2, the testing temperature237 in the range of 298 K to 673 K (25 C to 400 C) had a238 strong effect on alloy strength. The YS decreased from

244244244244244244161 to 84 MPa with the increasing testing temperature245from 298 K to 673 K (25 C to 400 C) (Figure 2(a)).246However, the YS remained about the same at ~145 MPa

(c)

(b)

(a)

#1

#2

#3#7

#4

#8

#9

#6

#10

#5

#5#8

#4#6

#3

#6#10

#5

#7

#4 #9

#1 #2

#3 #8

#5 #6

#7

#10

#3

#7

#7

#9

Fig. 1Microstructures of cast Al-Si-Cu-Mg alloy modified with Ti,V, and Zr, (a) OM image, (b) SEM image, and (c) SEM image afterdeep etching in BES mode.





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247between 473 K and 573 K (200 C and 300 C). In248contrast, the alloy UTS decreased linearly with the249testing temperature (Figure 2(b)). It is seen that the UTS250of the alloy decreased from 262 to 102 MPa with the251increasing temperature from 298 K to 673 K (25 C to252400 C) at a strain rate of 103 s1. It should be pointed253out that though the alloy modified with addition of Ti,254Zr, and V showed lower yield strength, the changing255trend in YS is lower than for the 354[7,27] and A356[37]

256alloys and the Al-Si-Cu-Mg alloy modified with addition257of Ti, Zr and Ni.[7,27] However, the studied alloy258achieved consistently higher UTS compared to the259published data for 354 (~12 pct), A356 (~18 pct), and260the Al-Si-Cu-Mg alloy modified with addition of Ti, Zr261and Ni in the as-cast state (~5 pct) at room temperature.262In comparison with the 354 and A356 alloys and the Al-263Si-Cu-Mg alloy modified with addition of Ti, Zr, and264Ni, the UTS at 473 K (200 C) for the present studied265alloy improved to ~11, ~94, and 7 pct, respectively. It is266believed that the presence of the thermally stable267precipitates improved the alloy strength at higher268temperatures[7,27] which will be discussed in later269sections.

TableII.

TheMainPhasesandTheirChem

istryinwtpctIdentified

UsingSEM/EDXintheCastAl-Si-Cu-M

gAlloyModified

withAdditionofZr-Ti-V

No.

CalculatedPhase

SuggestedPhase

Al

Si

Cu

Mg

Fe

Zr

Ti

V

#1

a-aluminum

a-aluminum

#2

eutecticsilicon

eutecticsilicon

35.282.31

64.722.31

#3

Al 2.1Cu

Al 2Cu

47.154.89

52.854.89

#4

Al 8.5Si 2.4Cu

Al-Al 2Cu-Si

63.556.54

18.774.54

17.681.96

#5

Al 7.2Si 8.3Cu2Mg6.9

Al 5Cu2Mg8Si 6

26.812.97

32.330.29

17.612.35

23.260.34

#6

Al 14Si 7.1FeM

g3.3

Al 8FeM

g3Si 6

52.842.06

28.031.43

11.281.39

7.851.58

#7

Al 21.4Si 4.1Ti 3.5VZr 3.9

(AlSi)3(TiVZr)

45.462.28

9.110.73

28.153.25

13.251.57

4.031.15

#8

Al 6.7Si 1.2TiZr 1.8

(AlSi)3(TiZr)

42.260.83

7.920.17

38.580.65

11.250.34

#9

Al 2.8Si 3.8V1.6Zr

(AlSi)2(VZr)

21.202.99

30.091.97

25.731.39

22.980.41

#10

Al 5.1Si 35.4Ti 1.6Zr 5.7Fe

(AlSi)2(TiZr)Fe

7.681.14

55.741.52

3.130.81

29.151.96

4.290.38

0

50

100

150

200

250

YS

, M

Pa

Testing temperature, K(C)

Mohamed et al., 2013 [27]

Present studied alloy


El-Kady et al., 2011 [37]

0

50

100

150

200

250

300

298(25) 473(200) 573(300) 673(400)

298(25) 473(200) 573(300) 673(400)

UT

S, M

Pa

Testing temperature, K(C)


Present studied alloy


El-Kady et al., 2011 [37]

(a)

(b) Fig. 2A comparison of the tensile property of the studied castAl-Si-Cu-Mg alloy with literature data obtained at different tem-peratures, (a) yield strength (YS) and (b) ultimate tensile strength(UTS). Note: the testing temperature for the study of Mohamedet al., 2013 was 463 K (190 C), instead of 473 K (200 C).





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270 2. Stress-strain behavior271 The true stress-true strain (r-e) plot of the studied272 alloy in the as-cast condition tested at room temperature273 is shown in Figure 3. The cyclic stress-total strain274 (ra Det/2) and the cyclic stress-plastic strain275 (ra Dep/2) curves of the as-cast alloy where ra is the276 stress amplitude at mid-life, Det is the total strain range,277 and Dep is the plastic strain range are shown in Figure 3.278 The obtained results show that the as-cast sample279 achieved a YS of 162 MPa, UTS of 261 MPa, and an280 elongation of 3.9 pct during monotonic tension loading281 as seen in Figure 3. It is evident that the studied alloy282 displayed significantly higher hardening behavior during283 cyclic loading compared to monotonic tensile loading. It284 is also remarkable that the alloy did not show any285 plastic deformation at lower strain amplitudes (0.1 to286 0.2 pct). However, with the increase in strain amplitude287 from 0.3 to 0.6 pct, the alloy exhibited noticeable plastic288 deformation. To better quantify the strain hardening289 behavior of the studied alloy, the stress-strain pa-290 rameters for uniaxial loading were characterized using291 the Hollomon power law,[23,36,38]

r Ken; 1

293293 where r is the true stress and e is the corresponding true294 strain. The strain hardening exponent n was evaluated295 for the uniform plastic deformation region between YS296 point and UTS point.297 The stress-strain parameters for cyclic loading were298 also characterized using the Hollomon power299 law,[23,36,38]

Dr

2 K0

Dep

2

n0; 2

301301 where Dr is the total stress range at mid-life and Dep is302 the corresponding total plastic strain range, n is the303 cyclic strain hardening exponent, and K is the cyclic304 strength coefficient. The as-cast samples exhibited a305 strain hardening exponent of 0.22 during monotonic306 tensile loading. In contrast, the n and K values of the

307presently investigated alloy were n = 0.09 and308K = 474 MPa, similar to the reference alloy modified309with the addition of Ti in amount of 0.1 to 0.14 wt pct.310(n = 0.115 to 0.154 and K = 335 to 380 MPa)[39] and311lower than the A356-T6 alloy (n = 0.24 and312K = 1628 MPa).[40] The alloy cyclic YS was 262 MPa313which was significantly higher than the monotonic314tensile loading. It is suggested that the alloys in the315present study had a stronger hardening ability under316cyclic loading than monotonic tensile loading. Indeed,317the evaluated cyclic strain hardening exponent318(n = 0.09) was clearly smaller than the monotonic319tensile strain hardening exponent (n = 0.22). Similar320results were reported by other researchers for monotonic321and cyclic deformation of Al-Si alloys.[28,36,40]

3223. Low-cycle fatigue properties323The plot of cyclic stress amplitude and the number of324cycles to failure at different strain amplitudes during325LCF tests is shown in Figure 4(a) for the studied alloy in326the as-cast condition. Generally, with the increasing327strain and stress amplitudes, the fatigue life decreased.

0

50

100

150

200

250

300

350

400

0 0.5 1 1.5 2 2.5 3 3.5 4

Stress,

MP

a

Strain, %

-

a- p/2 a- t/2

Fig. 3Stress-strain curves of cast Al-Si-Cu-Mg alloy with additionof Zr-Ti-V, obtained during tensile at a strain rate of 103 s1 andcyclic loading conditions at a strain rate of 102 s1.

0

50

100

150

200

250

300

Str

ess a

mp

litu

de, M

Pa

Number of cycles, N

0.10%

0.20%

0.30%

0.40%

0.50%

0.60%

(a)

00

.00

10

.00

20

.00

30

.00

4

1.E+0 1.E+1 1.E+2 1.E+3 1.E+4 1.E+5

1.E+0 1.E+1 1.E+2 1.E+3 1.E+4 1.E+5

Pla

sti

c s

train

am

plitu

de

Number of cycles, N

0.10%

0.20%

0.30%

0.40%

0.50%

0.60%

(b) Fig. 4Plot of the low-cycle fatigue tests data of cast Al-Si-Cu-Mgalloy with addition of Zr-Ti-V, (a) cyclic stress amplitudes vs thenumber of cycles and (b) cyclic plastic strain amplitudes vs the num-ber of cycles.





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328 The hardening behavior of the alloy can be observed in329 three groups: cyclically stable, initially attained harden-330 ing followed by reaching the saturation stage, and331 increasing hardening until failure. At lower total strain332 amplitude (0.1 pct) the cyclic stress amplitude remained333 basically constant throughout the entire LCF. At higher334 total strain amplitudes (0.4 to 0.6 pct), cyclic hardening335 occurred from the beginning and continued up to failure336 for the as-cast samples as cyclic deformation progressed.337 However, at 0.2 pct total strain amplitudes, initially the338 alloy showed cyclic hardening up to 1000 cycles, after339 that the sample reached cyclic stability. After reaching340 the saturation level, cyclic stability was observed during341 cyclic deformation. The tendency toward cyclic harden-342 ing became stronger with the increasing strain ampli-343 tudes from 0.3 to 0.6 pct, as indicated by the increasing344 slope in the semi-log scale diagram (Figure 4(a)). The345 cyclic deformation features occurring in the present346 alloy were observed to be in agreement with those347 reported in References 28, 29, 40 for Al-Si alloys.348 Another notable change was observed in the curves of349 cyclic stress amplitude as a function of the number of350 cycles at different total strain amplitudes, i.e., the values351 of initial cyclic stress amplitudes increased from ~70 to352 ~195 MPa with the increase in total strain amplitude353 from 0.1 to 0.6 pct, respectively. The alloy cyclic354 hardening behavior depends on the yield strength of355 the material. Cyclic hardening and plastic deformation356 occurred, if the total stress amplitude is higher than the357 yield strength of the alloy. At lower strain amplitude358 (0.1 pct), the yield strength (161 MPa) was higher than359 the maximum total stress amplitude (~70 MPa), while at360 higher total strain amplitudes (0.2 to 0.6 pct), the361 maximum total stress amplitude was much closer362 (~125 to ~195 MPa) to the yield strength (161 MPa) of363 the studied alloy. This resulted in a lack of hardening at364 lower strain amplitudes of the studied alloy, while at365 higher strain amplitudes the alloy underwent hardening366 throughout its entire fatigue life.367 Figure 4(b) shows the plastic strain amplitude as a368 function of the number of cycles for the studied alloy.369 The fatigue life of the alloy decreases with the increase in370 total strain amplitude leading to an increase in the371 plastic strain amplitude. A linear decrement of plastic372 strain amplitude was observed at higher total strain373 amplitudes (0.4 to 0.6 pct), while it remained almost374 constant at lower total strain amplitude of 0.1 pct.375 However, at lower total strain amplitudes of 0.2 to376 0.3 pct, initially the alloy plastic strain amplitude377 decreased followed by constant plastic strain amplitude378 for the remaining fatigue life. This corresponded well to379 the cyclic hardening and stability characteristics seen in380 Figure 4(a) for the same total strain amplitude.381 Figure 5 shows the plot of strain and the number of382 cycles to failure (generally known as e-N curves) for the383 studied alloy in the as-cast condition along with a384 reference A356 alloy in the as-cast state. It is seen that385 the fatigue life increased with the decrease in total strain386 amplitudes for all the alloys. The fatigue life of the387 presently developed alloy appeared slightly longer than388 the A356 alloy[41] and other existing reference T6389 tempered Al-Si-Cu-Mg alloys modified with additions

390of Zr, Ti, and V at all levels of total strain amplitudes.[28]

391It should be noted that if the alloy fatigue life reached392107 cycles without failure, fatigue test was discontinued.393Hence, at the strain amplitude of 0.1 pct, the present394studied alloy passed the fatigue test, while the existing395Al-Si-Cu-Mg alloy modified with the addition of Zr, Ti,396and V failed to achieve an infinite fatigue life.397During LCF, the total strain amplitude consists of398elastic strain amplitude and plastic strain amplitude as399shown in the following relation[28,29,36]:

Det

2Dee

2Dep

2; 3

401401where Det2is the total elastic strain amplitude, Dee

2is the

402elastic strain amplitude andDep2

is the plastic strain403amplitude.404The fatigue parameters, Nf is the number of cycles to405failure, r0f is the fatigue strength coefficient, b is the406fatigue strength exponent, e0f is the fatigue ductility407coefficient, and c is the fatigue ductility exponent, are408described in the Basquin equation, the Coffin-Manson409equation, and the Coffin-Manson-Basquin equation in410References 29, 36. The values of the fatigue parameters411were evaluated using linear regression analysis as412described in References 29. The calculated fatigue413parameters are listed in Table III for the present studied414alloy. It is seen that the alloy had a fatigue strength415coefficient of 386 MPa and the fatigue strength expo-416nent of 0.08. Also, the alloy showed a fatigue ductility

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

1E+0 1E+2 1E+4 1E+6 1E+8

t/2, %

Number of cycles to failure, Nf

Elhadari et al., 2011 [28]



Azadi and Shirazabad,2013 [41]Present study alloy

Fig. 5Fatigue life of the cast Al-Si-Cu-Mg alloy with addition ofTi, V, and Zr obtained at total strain amplitudes of 0.1 to 0.6 pct incomparison with the existing reference alloys.

Table III. Evaluated Materials Constants for LCF of theCast Al-Si-Cu-Mg Alloy with Addition of Ti, V, and Zr in

Different Amounts

Cyclic yield Strength, r0y (MPa) 262Cyclic strain hardening exponent, n 0.09Cyclic strength coefficient, K (MPa) 474Fatigue strength coefficient, r0f (MPa) 386Fatigue strength exponent, b 0.08Fatigue ductility coefficient, e0f (pct) 9.21Fatigue ductility exponent, c 0.77





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417 coefficient and exponent of 9.21 pct and 0.77, respec-418 tively, which were significantly higher in the present419 investigated alloy in the as-cast condition than the420 existing Al-Si-Cu-Mg alloy modified with minor addi-421 tions of Zr, Ti, and V in the T6 state.[28] Thus, it could422 be concluded that the present alloy has slightly better423 fatigue properties compared to the previously tested T6424 tempered Al-Si-Cu-Mg alloy modified with minor addi-425 tion of Zr, Ti, and V.[28]

426 C. Fractography

427 1. Analysis of fracture surfaces after tensile tests428 The overview of specimens of the studied alloy429 fractured after tension tests at temperatures of 298 K,430 473 K, 573 K, and 673 K (25 C, 200 C, 300 C, and431 400 C) is depicted in Figure 6. Generally, the ductile432 fracture occurred by formation, accumulation, and433 growth of voids due to cracking of the second-phase434 precipitates. The voids were formed by raising the435 interfacial stress which in turn resulted in the breaking436 of the interfacial bonds, due to critical stress, between437 the precipitate and the ductile matrix. An alternative438 mechanism is an initiation of cracks within hard439 precipitates. However, casting defects, especially inter-440 nal pores, dry oxide, or shrinkage etc., act as preferable441 sites for crack nucleation as well.

442The fractographic observations revealed that the443fracture characteristics were changed from mixed mode444to ductile fracture with the increase in testing tem-445perature from 298 K to 673 K (25 C to 400 C)446(Figures 7 and 8). After tensile testing at room tem-447perature [298 K (25 C)], the fracture surface exhibited448mostly intergranular features along with some flat areas,449especially cleavage-type fracture in the plate-shaped450intermetallics (Figures 6(a) and 7(a)). Some eutectic451silicon precipitates were debonded from the matrix (as452pointed out by red arrow) which was accompanied by453secondary cracks between dendrites (blue arrows)454(Figure 8(a)). Such morphology suggests that there455was a strong interaction between the plastic flow or slip456bands and the eutectic silicon precipitates especially at457grain boundaries leading to a contribution to inter-458granular cracking. It is also obvious that there is459microporosity (enclosed by white-dashed line) on the460fracture surface of samples tested at room temperature461and higher temperatures as well (Figure 6). Similar types462of fracture behavior were also pointed out by other463researchers for the Al-Si-Cu-Mg alloy modified with the464addition of Zr, Ti, and V during tensile loading.[17,28]

465A different pattern of fracture is observed in the466specimen tested at 473 K (200 C). The flatness of the467fracture surface in Figure 6(b), coupled with reduced468volume fraction of precipitates, suggests a diminished

(c)

(b)(a)

(d)

#m#m

#m#m

#p

#p

#p #p

Fig. 6SEM micrographs showing the overall tensile fracture surface of studied alloy in as-cast states obtained at different temperatures of (a)298 K (25 C), (b) 473 K (200 C), (c) 573 K (300 C), and (d) 673 K (400 C). Note: enclosed p and m areas are magnified in Figs. 7 and 8,respectively.





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469 role of precipitates on fracture at this temperature. As470 reported in Reference 42, the relative weakness of the471 grain boundary compared to the room temperature472 could be a possible reason for different fracture473 mechanism. Fracture on a microscopic scale comprised474 cracking of the clusters of Si precipitates and the Zr-Ti-475 V-rich intermetallics present in the microstructure. It is476 observed that the Zr-Ti-V-rich intermetallics were pre-477 sent on several cleavage planes (Figure 7(b)). The brittle478 crystalline fracture with small areas of plastic deforma-479 tion is noticed on the fracture surface of the specimen480 tested at 473 K (200 C) (Figures 7(b) and 8(b)). Here,481 brittle cleavage fracture is dominating up to 473 K482 (200 C). Above 473 K (200 C), i.e., at 573 K (300 C),483 the fracture surface indicates the brittle-to-ductile tran-484 sition due to the effect of temperature. In this circum-485 stance, the Zr-Ti-V-rich intermetallics also played an

486effective roll in fracture behavior of the alloys as seen487from the EDS analysis in Figure 7. At this higher488temperature, interfacial bonding between the matrix and489Zr-Ti-V-rich intermetallics became weaker which en-490hanced decohesion of the precipitates from the matrix491resulting in occasionally pull-out of the Zr-Ti-V-rich492intermetallics (Figure 7(c)). This phenomenon is more493frequent for the specimen tested at 673 K (400 C) as494enclosed by red line (Figure 7(d)). It can be attributed to495the different modulus and thermal expansion of the496precipitates (phase #7 in Table II) and matrix, softening497of the matrix, and partial dissolution of the precipitates498(phases # 8 and 9 in Table II) in the matrix. When the499temperature rises, the thermal expansion of the matrix500and the precipitates creates a gradient that increased the501interfacial stress leading to fracture and pull-out of the502precipitates. Also, the microscopic observation at high

(c)

(b)(a)

(d)Fig. 7SEM micrographs with EDX spectra in a magnified view correspondingly showing the tensile fracture surface (enclosed in Fig. 6 as p) ofstudied alloy in as-cast states obtained at different temperatures of (a) 298 K (25 C), (b) 473 K (200 C), (c) 573 K (300 C), and (d) 673 K(400 C). Note: the areas enclose by white and red lines show the surface pores and pull-out of the precipitates, respectively.





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503 magnifications revealed that the fracture surface con-504 sisted of a population of micro-voids for samples tested505 at 573 K (300 C) (Figure 8(c)). The nucleation and506 coalescence of the dimples are not visible and the alloy507 will flow until cracking is complete. The homogeneous508 distribution of micro-voids with different size range was509 observed in this condition (Figure 8(c)). When the510 temperature increased to 673 K (400 C), the matrix511 became much softer and localized plastic deformation512 occurred which accumulated the voids and enlarged the513 cavity (Figure 8(d)). Thus, the fracture surface at 673 K514 (400 C) is dominated by ductile trans-crystalline frac-515 ture in the studied alloy.516 Figure 9 illustrates the SEM observations of the517 polished cross sections of the studied alloy after tensile518 testing at different temperatures. As seen in Figures 9(a)519 and (b), the tensile test sample shows that secondary520 cracking and small amounts of plastic deformation521 occurred near the fracture surface during tests at room522 temperature and 473 K (200 C). In contrast, with the523 increase in the testing temperature above 473 K524 (200 C), plastic deformation of the alloy increased525 and its fracture mode changed from mixed fracture to526 ductile fracture (Figures 9(c) and (d)). At the same time,527 several of the secondary cracks are the result of528 accumulation of voids generated due to cracking of529 precipitates. As seen in Figure 10, the micro-cracking530 which is generally parallel to the loading axis occurred in531 the intermetallics (indicated by red arrows) and eutectic

532silicon (marked by blue arrows). Multiple cracks were533found mainly in the high-aspect ratio precipitates, due534to strong interfacial bonding between the precipitate and535the aluminum matrix. Occasionally, debonding between536the precipitate and matrix was also seen after tensile537tests at room temperature and at 473 K (200 C). In538contrast, after tensile testing at 573 K (300 C), addi-539tional multiple micro-cracking phenomena in the inter-540metallics were recorded. Those cracked precipitates are541essentially aligning themselves along the loading axis (as542enclosed by green line in Figure 10(c)). Micro-cracks in543precipitates were also revealed in the sample tested at544673 K (400 C). At the same time, precipitates are545displaced from their original position and aligned to the546loading direction (as enclosed by green line in547Figure 10(d)). This is an indication of weak interface548which enhances decohesion and pull-out of the pre-549cipitates at higher temperatures.

5502. Analysis of fracture surfaces after LCF551The overall features of the fatigue fracture surfaces552are presented by SEM images in Figure 11. Here, the553crack initiation site and propagation zone for the554studied alloy tested at 0.2 and 0.6 pct of the total strain555amplitude are illustrated. As seen in Figure 11(a), there556are three distinct zones observed on the fracture surface;557i.e., (i) fatigue crack initiation (FCI) zone, (ii) fatigue558crack growth/propagation (FCG) zone, and (iii) final559fracture zone which is adjacent to the uneven FCG zone.

(c)

(b)(a)

(d)Fig. 8The corresponding SEM micrographs showing the matrix morphology (enclosed in Fig. 6 as m) of the tensile fracture surface of studiedalloy in as-cast states obtained at different temperatures of (a) 298 K (25 C), (b) 473 K (200 C), (c) 573 K (300 C), and (d) 673 K (400 C).





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560 It is noticed that with the increasing total strain561 amplitude, the size of the propagation zone decreases562 (indicated by dashed red line in Figure 11(a)). However,563 the fracture surface of the sample tested at a higher564 strain amplitude (0.6 pct) shows negligible crack565 propagation zone, i.e., resembling room temperature566 tensile-like fracture behavior (as discussed in the previ-567 ous Section IIIC1) due to fast failure (Figure 11(b)).568 The higher magnification images depicted in569 Figures 11(c) and (d) showed that crack initiation oc-570 curred from intrinsic slip band (as marked by red arrow571 the small half circle flat surface) and casting defects near572 the sample surface, such as a large surface pore (as573 enclosed by white line) for the sample tested at a total574 strain amplitude of 0.2 pct in Figure 11(c), while a large575 near-surface pore initiated the cracks for the sample tested576 at the total strain amplitude of 0.6 pct in Figure 11(d).577 These locations acted as stress concentrators that initiated578 the formation of fatigue cracks.[40,43,44]

579 The area enclosed by white boxes on the fracture580 surface in Figure 11(a) is further magnified and pre-581 sented in Figures 12(a) through (f), where more detailed582 features in the fatigue propagation and fracture zones of583 the studied alloy, tested at a total strain amplitude of584 0.2 pct, could be observed. It is clearly seen that the585 crack propagation zone exhibited fine fatigue striations,586 which are surrounded by tear ridges of the grain587 boundaries, which are apparent on the fracture surfaces588 of the sample tested at a lower strain amplitude (0.2 pct)

589in Figures 12(a) and (b). Another interesting feature590recognized on the fracture surface was the prominent591micro-cliff on which fatigue striation was formed along592the crack propagation plane (Figure 12(a)). These mi-593cro-cliffs have a step-like pattern formed inside the grain594generally parallel to the FCG path, which seem to595indicate lateral slippage at the crack tip.[36,45] As596portrayed in Figures 12(c) and (d), the fatigue striations597were also noticed on the Zr-Ti-V-containing precipitates598of the sample tested at the lower strain amplitude of5990.2 pct. As noticed in Figure 12(e), the EDX analysis of600the enclosed area in Figure 12(c) confirms the presence601of Zr-Ti-V-containing precipitates on the fracture sur-602face. It is obvious that multiple fatigue striations603morphology, as indicated by dotted line in Figure 12(d),604were initiated and then overlapped with each other is an605indication of different set of crack propagation paths606through particles, as shown by red-dotted box in607Figure 12(d). The fatigue striations are only possible in608ductile materials. Thus, it can be presumed that the Zr-609Ti-V-containing precipitates are relatively ductile com-610pared to other intermetallics present in this study, which611essentially improved the LCF performance of the alloy.612In the final fracture zone (Figure 12(f)), the above-613discussed tensile features are observed. Also, some614secondary cracks in the matrix were detected in the615stage of the final fracture (Figure 12(f)). Similar features616of LCF fracture characteristics were also observed by617other researcher in cast aluminum alloys.[28,36,40] Also,

(a) (b)

(d)(c)

Fig. 9SEM images illustrating the overall view of the polished cross section of the tensile-tested samples near the fracture surface, obtained atdifferent temperatures of (a) 298 K (25 C), (b) 473 K (200 C), (c) 573 K (300 C), and (d) 673 K (400 C). Note: enclosed areas are magnifiedin Fig. 10.





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618 more details about the development of fatigue striation619 in the matrix and Zr-Ti-V-rich intermetallics are620 documented in our recent publication.[29]

621 IV. DISCUSSION

622 The cast Al-Si-Cu-Mg alloy microstructure depends623 on the modification of the eutectic silicon, solidification624 rate, heat treatment processes, as well as alloy chemistry.625 In other words, the processing and compositional626 parameters considerably influence the tensile and fatigue627 properties of the cast Al-Si-Cu-Mg alloys. The presented628 results are explained in the following sections by629 addressing the effects of testing temperatures, yield630 strength, and the behavior of precipitates, especially Zr-631 V-Ti-rich phase on the alloy performance.

632 A. Influence of Temperature on the Alloy Strength

633 Present results reveal that Al-Si-Cu-Mg modified634 alloys containing Zr, V, and Ti exhibited better tensile635 strength compared to reference alloys,[27] at different636 temperatures as seen in Figure 2. It is noticed that the637 alloy room temperature tensile strength is significantly638 higher than the alloy modified with addition of Zr-Ni in639 alloy 354. As mentioned by Mohamed et al.,[27] addition

640of Ni to the Al-Si-Cu-Mg alloy has a poisonous effect on641the age hardening which reduces the copper content in642the matrix and lowers the strength of the alloy. In643contrast, the addition of Zr-V-Ti to the Al-Si-Cu-Mg644alloy caused thermally stable precipitates to form during645solidification of the alloy in conjunction with the646copper-containing phases that together increased the647alloy strength. Also, the detrimental Fe-containing648phase reacts with Zr-Ti and forms thermally stable649precipitates which enhance the alloy performance. As650reported in previous studies,[28,40] the remarkable651enhancement in UTS could be attributed to both the652composite-like role of Cu-, Fe-, and Mg-containing Si653precipitates (Figure 1), and the nano-sized trialuminide654precipitates which were uniformly distributed in the655aluminum matrix or along with the eutectic silicon656precipitates. The presence of these micro- and nano-657sized precipitates (Figure 1) would effectively impede658the movement of dislocations during uniaxial deforma-659tion, thus appreciably enhancing the strength of the cast660aluminum alloy (Figure 2). Also, the addition of tran-661sition metals containing Zr-V-Ti to Al alloys changes662the morphology of the primary precipitates which can663improve the strength and ductility of the alloys.[16]

664Therefore, it is reasonable to say that the bulk-/plate-665shaped Zr-Ti-V-rich phases improved the as-cast alloy666strength.

(b)(a)

(d)(c)

Ten

sion

axis

Micro-voids

Micro-voids

Micro-voids

Fig. 10SEM images illustrating the magnified view (enclosed in Fig. 9 by red box) of the polished cross section of the tensile-tested samples,acquired at different temperatures of (a) 298 K (25 C), (b) 473 K (200 C), (c) 573 K (300 C), and (d) 673 K (400 C). Note: the areas enclosedby green lines show the displacement of the precipitates.





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667 As mentioned above, with the increasing testing668 temperature, the tensile strength of Al-Si-Cu-Mg alloy669 modified with addition of Zr, Ti, and V decreased670 linearly. This was primarily due to the change in the671 microstructure and the presence of trialuminide pre-672 cipitates, as also reported in References 14, 16. The673 deformation in crystalline materials is directly related to674 the generation, motion, and storage of dislocations,675 which move at a stress far below that required to deform676 a defect-free crystal.[46] Therefore, the improvement of677 material properties such as strength relies on creating678 internal obstacles to dislocation motion including grain679 boundaries, precipitates, and other dislocations.[47,48] As680 plastic deformation continues, the density of disloca-681 tions increases, causing more frequent interactions that682 impede their motion.[49] When the temperature is683 increased, the cross-slips are thermally activated by684 climbing of dislocations resulting in reduction of685 strength of the materials.[7,27] At the same time, grain686 boundary sliding is also considered to be a reason for687 the reduction of the alloy strength at higher tem-688 peratures. However, the strong resistance to the motion689 of dislocations, due to the trialuminide precipitates690 pinning grain and sub-grain boundaries during all

691thermal and mechanical processing of aluminum alloys,692causes the improvement of the high-temperature prop-693erties of the alloy.[27,28,35,40,50,51]

694Another fascinating phenomenon was observed for695the behavior of intermetallics present on the fracture696surfaces of the studied alloy. As seen in Figures 7 and6978, mixed-type (brittle particle and ductile matrix)698fracture dominated at lower temperature, while duc-699tile-type fracture was identified at higher temperatures.700The brittle fracture may be dominated by silicon701particles and intermetallics, especially Zr-V-Ti-rich702phases. Brittle cleavage and mixed fracture mode with703localized plastic deformation is governing up to 473 K704(200 C). Beyond 473 K (200 C), the fracture surface705indicates the brittle-to-ductile transition due to the706effect of temperature. As discussed by Wang,[52] the707fracture mechanism of the materials can be divided708into three steps: (i) particle cracking which depends on709the localized condition including particle size, particle710shape, particle orientation, activation of dislocation711source, etc., (ii) generation and growth of micro-voids,712and (iii) linkage of micro-voids followed by final713fracture. As schematically represented in Figure 13,714the Zr-V-Ti-rich precipitates cracked during tensile

(c)

(b)(a)

(d)

FCI

FCI

#i #ii

Porosity

Porosity

FCG

FCG

#iii

Fig. 11SEM images of fatigue fracture surfaces of the samples tested at a total strain amplitude of (a) 0.2 pct and (b) 0.6 pct showing an over-all view for cast Al-Si-Cu-Mg alloy modified with addition of Zr-Ti-V and the corresponding fatigue crack initiation sites at higher magnifica-tions, where fatigue cracks initiate at a large surface pore and slip bands (c); clusters of near-surface pores and large surface inclusion (d). Here,yellow box and red arrows indicate the position of crack initiation sites, while red dashed line separated the crack propagation area. The areasenclose by white line shows the surface pores. Note that increasing the strain amplitude crack propagation zone decreases. FCI: fatigue crackinitiation, FCG: fatigue crack growth/propagation zone.





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715 loading. Due to particle cracking, micro-voids are716 generated around the particles. Those micro-voids are717 connecting with each other, resulting in final fracture.718 During tensile loading, particles experience multiple719 cracking without any debonding, which is an indication720 of the strong interfacial bonding between the particles721 and Al-matrix at the testing temperature up to 473 K722 (200 C) (Figure 13(b)). However, the fractured parti-723 cles are debonded from the matrix and pull-out from724 the sub-surfaces for the samples deformed at higher725 temperature (Figures 7(d) and 13(c)). This can be726 attributed by the fact that softening of the solid occurs727 with the increase in temperature beyond 0.7Tm (melting728 temperature) at which the solid loses the general729 properties.[53] At the same time, the dislocations are

730piled up at the interface of the particles and matrix731resulting in debonding and pull-out. Further, the732ductile fracture is determined by the size of dimples,733being governed by the number and distribution of734micro-voids that are nucleated. The microscopic ob-735servation at high magnification reveals that the fracture736surface consists of a population of micro-voids (Fig-737ures 8(c) and (d)). The sources of voids and resultant738dimples on the fracture surface are attributed to the739fracture of the Zr-V-Ti intermetallic particles. Finally,740it can be concluded that the fracture of the studied741alloy at temperature 673 K (400 C) is related to the742void nucleation and growth. The nucleation is initiated743by the coarse constituent particles and other second744phases present in the alloy microstructure.

(c)

(b)

(d)

(a)

Element Wt.%

Al 48.22

Si 9.66

Ti 9.79

V 2.59

Zr 29.74

(e)

Dimples

Micro-cliffs

Tear ridge

(f)

Tear ridge

Fatigue striation

Fatigue striation

Fig. 12SEM images of fatigue fracture surfaces for cast Al-Si-Cu-Mg alloy modified with addition of Zr-Ti-V samples tested at total strainamplitude of 0.2 pct showing fatigue striation in matrix (a and b) (as indicated in Fig. 11(a) with i) and precipitates (c and d), (as pointed out inFig. 11(a) with ii), (e) EDX analysis of intermetallic (typical #7 as listed in Table II) on the fracture surface and (f) final fracture of the alloy (asmarked in Fig. 11(a) with iii). Note: the dotted lines in (e) are two sets of fatigue striation along the crack propagation direction overlappedeach other as indicated by enclosed red box.





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745 B. Improvement of LCF Performance

746 The LCF results also showed superior properties over747 the existing reference alloys. The improved fatigue life of748 the present studied alloy was mainly attributed to the749 thermally stable trialuminide phases which significantly750 increased precipitate volume fraction and changed751 morphology upon introducing more alloying elements752 (i.e., Zr, Ti, and V). The obtained cyclic yield strength of753 the studied alloy was always higher than the monotonic754 yield strengths which indicate the higher hardening and755 better dislocation storage capacity in cyclic loading. As756 discussed earlier, the hardening behavior of the alloy can757 be explained by Hollomon parameters (Eq. [2]). Gener-758 ally, lower K and n denote the faster velocity of cyclic759 hardening, higher plastic deformation resistance, and760 lower ductility. It was reported that the cyclic softening761 appears at n< 0.1, cyclic stability at n = 0.1, and762 cyclic hardening at n> 0.1.[39] The value of n = 0.09763 for the studied alloy is greater than 0.1, confirming

764cyclic hardening of the alloy. The values of765K = 474 MPa and n are lower than those for the766reference alloy (n = 0.28 and K = 2393 MPa) and the767A356-T6 alloy (n = 0.24 and K = 1628 MPa),[40]

768indicative of a higher hardening rate for the studied769alloy.770For fatigue, a tension-compression process accompa-771nies the whole failure. It is speculated that there exists a772fixed forward and backward movement for some mobile773dislocations at this stage. With the increasing number of774fatigue cycles, the dislocation density gradually increases775and the dislocation entanglement becomes more disor-776derly, which enhances the resistance of subsequent777dislocation movement and gives rise to the evident778cyclic hardening trend. Also, the effect of the solidifica-779tion microstructure on the cyclic hardening behavior is780mainly attributed to the interaction between the dislo-781cation and the second phases (eutectic silicon and782intermetallics) or the grain boundaries.[39,54] When the

(b)

(c)

(a) #7/8/9

Fig. 13A schematic model depicting the successive steps of fracture mechanism of the studied alloy at room temperature (a and b) and hightemperature (a and c) during tensile loading. The #7/8/9 are the precipitates listed in Table II. Note: during tensile loading micro-cracks whichformed micro-voids in matrix were formed. The micro-voids connect with each other leading to the ultimate failure of the alloy.





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783 cyclic deformation under alternating loading takes784 place, the cross-slip ability of dislocation will con-785 tinuously increase.[39,55] The resistance of dislocation786 slipping comes from the boundaries between the Si787 phase and the grains, and the refiner grains consequen-788 tially shorten the mean free path of dislocation slipping789 and thus result in the higher cyclic hardening ability. As790 discussed earlier, the LCF performance of materials is791 also indicated by the parameters in the equation of792 Coffin-Manson-Basquin.[56] Higher values of b and c793 indicate higher cyclic strength and ductility of the794 materials. The present studied alloy shows a higher795 value of c, compared to the existing reference alloy in796 the T6 heat-treated condition, indicative of higher797 ductility of the alloy; leading to longer fatigue life.798 Tensile property also influences the LCF cyclic799 hardening behavior of metallic materials, indicative of800 higher hardening and better dislocation storage capacity801 in cyclic loading. If the monotonic YS is lower and the802 UTS is higher, the size of the plastic zone at the crack803 will be larger, which will blunt and limit cracks from804 propagating by improving the matrix hardening. On the805 other hand, the movement of dislocations from cell or806 grain boundaries is inhibited which reduces the interac-807 tion between the dislocations and precipitates resulting808 in the reduction of crack propagation and improvement809 of fatigue performance. The studied alloy obtained810 comparatively lower YS and consistently higher UTS811 compared to the reference alloys (Figure 2). At the same812 time, the LCF life also depends on the alloy microstruc-813 ture. The studied alloy consists of a large number of814 alloying elements, which develop a complicated mi-815 crostructure by forming intermetallic and eutectic phas-816 es (listed in Table II). Those intermetallics in817 conjunction with the eutectic Si precipitates separate818 the aluminum matrix into a large number of tiny819 domains (as depicted in Figure 1). During LCF, the820 dislocations are moving into those domains of the821 aluminum matrix. Thus, the surrounding particles resist822 the slip bands or strain localization in the aluminum823 matrix which increases the LCF life of the alloy. Thus,824 the LCF of the studied alloy is improved slightly.825 As seen in Figures 11 and 12, the fatigue crack826 preferred to propagate along the boundaries between827 the intermetallics and Al-Si eutectics due to debonding828 or fracture of intermetallics and Si particles. It is also829 noted that the fatigue cracks propagated through the830 plate-shaped particles especially Zr-Ti-V-rich inter-831 metallics. It can be assumed that the stiff intermetallics832 could not follow the flexible Al matrix to deform and833 thus micro-cracks form at the interface of intermetallics.834 These micro-cracks will bring out micro-voids in the835 boundaries. These micro-voids in intermetallics or836 eutectics accumulate and propagate along their bound-837 aries and finally connect each other to weaken the838 boundaries.[57] On the contrary, EDS analysis shows839 that most of the facets on the tensile/fatigue fracture840 surface are the Zr-Ti-V-rich intermetallics (Fig-841 ure 12(e)). When the tensile crack propagates in this842 Al-Si-Cu-Mg alloy, the intermetallics with different843 orientations are actually obstacles to crack growth.844 Suffering from the strong crack-tip stress, these Zr-Ti-V-

845rich intermetallics are fractured ahead of the crack tip846without obvious debonding. Thus, the presence of Zr-847Ti-V-rich intermetallics slightly improved the fatigue848performance of the studied alloy.

849V. CONCLUSIONS

850To improve the high-temperature tensile properties851and low-cycle fatigue performance, the Al-Si-Cu-Mg-852base alloy was modified with the addition of Ti, V, and853Zr. From the above-presented results, the following854conclusion can be drawn:855The Al-Si-Cu-Mg alloy studied in this work and856modified with addition of Zr, Ti, and V developed a857complex microstructure. The EDX analysis of the Zr-Ti-858V-rich phases showed that Zr always form intermetallics859with other transition metals (Ti and V) and with the860increasing Zr content in the Zr-Ti-V-rich phases, Si and861V content decreases. The developed micro-sized Zr-Ti-862V-rich phases in the studied alloy, Al21.4Si4.1Ti3.5VZr3.9,863Al6.7Si1.2TiZr1.8, Al2.8Si3.8V1.6Zr, and Al5.1Si35.4-864Ti1.6Zr5.7Fe were calculated from the EDX data.865The tensile tests at different temperatures showed that866the addition of Zr, Ti, and V to the Al-Si-Cu-Mg alloy867successfully improved the alloy strength, which had868significantly higher UTS in comparison with the A356869and 354, and the existing Al-Si-Cu-Mg alloy modified870with the addition of Zr, Ti, and Ni. It is also noted that871increasing the testing temperature from 298 K to 673 K872(25 C to 400 C) of the present studied alloy decreased873the YS and UTS from 161 to 84 MPa and from 261 to874102 MPa, respectively.875The fatigue life of the studied alloy was considerably876longer than that of the reference alloy A356 and the alloy877with the same base but minor addition of V, Zr, and Ti in878the T6 state, reported in the literature. The cyclic stress879amplitude and plastic strain amplitude were almost880stable at low total strain amplitude of 0.1 pct. Consis-881tently, cyclic hardening occurred at higher strain ampli-882tudes with the extent of cyclic hardening increasing with883the increasing total strain amplitudes from 0.3 to 0.6 pct.884Fractographic analysis of tensile-tested samples885showed that mixed-type (cleavage-type fracture and886brittle fracture of particles and ductile fracture of the887matrix) fracture dominated at lower temperature while888ductile-type fracture was identified at higher tem-889perature. Multiple cracks were observed in the Zr-Ti-890V-rich phases indicating the strong interfacial bonding891between the precipitates and the matrix at temperature892up to 473 K (200 C). However, after increasing the893testing temperature from 473 K to 673 K (200 C to894400 C), debonding and pull-out of cracked inter-895metallics were identified on the alloy fracture surface.

896

897ACKNOWLEDGMENT

898The authors would like to acknowledge the financial899support of the ecoENERGY Innovation Initiative900ecoEII of Natural Resources Canada at Can-





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901 metMATERIALS. One of the authors (D.L. Chen) is902 grateful for the financial support by the Natural Sci-903 ences and Engineering Research Council of Canada904 (NSERC), PREA, NSERC-DAS Award, CFI, and905 RRC program. The authors would also like to thank906 Q. Li, A. Machin, J. Amankrah, and R. Churaman907 for their assistance in the experiments. The authors908 also thank Professor S. Bhole for the helpful discus-909 sion as well as P. Newcombe, G. Birsan, H. Webster,910 D. McFarlan, F. Benkel, M. Thomas, and D. Saleh911 from CanmetMATERIALS for casting experiments.

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