New heat treatment to improve the mechanical properties of low copper aluminum primary foundry alloy

Accepted Manuscript New heat treatment to improve the mechanical properties of low copper aluminum primary foundry alloy R. Martinez, I. Guillot, D. Massinon PII:

S0921-5093(19)30447-2

DOI:

https://doi.org/10.1016/j.msea.2019.04.001

Reference:

MSA 37751

To appear in:

Materials Science & Engineering A

Received Date: 23 December 2018 Revised Date:

30 March 2019

Accepted Date: 1 April 2019

Please cite this article as: R. Martinez, I. Guillot, D. Massinon, New heat treatment to improve the mechanical properties of low copper aluminum primary foundry alloy, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/j.msea.2019.04.001. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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New heat treatment to improve the mechanical properties of low copper aluminum primary foundry alloy R. Martineza , I. Guillotb,∗, D. Massinonc a

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LINAMAR CORPORATION, 700 Woodlawn Rd West, Guelph, Ontario, Canada, N1K 1G4 b Université Paris-Est, ICMPE (UMR 7182), CNRS, UPEC, F-94320 Thiais, France c LINAMAR MONTUPET Light Metal Casting Division, 3 rue de Nogent, 60290 Laigneville, France

Abstract

The purpose of this work is to design a new heat treatment in order to improve the mechanical properties of primary 356+0.5wt%Cu aluminum alloy. This heat treatment combines a multi step solution treatment with a quench and

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a multi step overaging treatment. It takes advantage of the evolution of the solubility limit of the matrix as a function of temperature in order to allow the appearance of successive precipitation sequences during excessive aging in

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several stages and leads to an increase in the volume fraction of precipitates. After T64 air quench, T6 air quench, T7 air quench and T7 water quench

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treatments and overaging 100h at 150◦ C, a secondary precipitation of Al2 Cu particles is seen using conventional TEM, mostly after the T7 treatments. This secondary precipitation leads to an increase of the static mechanical properties. This dual-step overaging is then transposed to a dual-step aging T7 heat treatment, with aging 2h at 200◦ C and 2h at 170◦ C, and further compared to a classic T7 treatment with aging during 4h at 200◦ C. Both heat 1

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treatments are also modeled using a KWN model. The model captures the secondary supersaturation of the proeutectic α-phase that leads to the sec-

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ondary precipitation during the dual-step aging treatment. This translates to

an increase of the mechanical properties by 6% and 3% for the yield strength

and the ultimate tensile strength respectively. The elongation decreases by

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33%.

Keywords: Heat treatment, step-aging, secondary precipitation, foundry

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aluminum alloys, cylinder heads, Transmission Electron Microscopy.

1. Introduction

For the past twenty five years, weight reduction has been the main motivation for using aluminum alloys in the automobile industry. At first, only

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massive parts manufactured in cast iron made the switch to aluminum alloys. Indeed, these alloys have excellent specific mechanical/thermal properties and they are easily recyclable which make them an excellent alternative to

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ferrous alloys. As a consequence, cylinder heads were amongst the first

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parts to embrace lightweighting.

Secondary 319 alloys (3wt.% Cu) have been successfully used for cylin-

der heads application until very recently. The possibility to use the alloy without heat treatment (also known as F temper) or with a simplified T5 heat treatment added to the cost benefit of using a secondary alloy, made it suitable ∗

Tel.: +33 (0)1 56 70 30 55 Email address: [email protected] (I. Guillot) Preprint submitted to Elsevier

April 4, 2019

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for low to medium performance engine cylinder heads. However, because of poor damage tolerance, low heat conductivity and issues with residual stress

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when conducting a full T6 heat treatment, 319 secondary alloys were slowly abandoned.

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More recently, T7 heat treated primary low copper 356 and AlSi10 alloys have gained popularity for manufacturing highly loaded cylinder heads. In-

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deed, associated to tilted gravity process they offer the best compromise to cast heavily loaded and complex cylinder heads with multiple cores and integrated exhaust manifolds. There low amount of impurities (Fe, Ni, Cr, Na and Ca) and there low level of copper (typically ≈0.5wt.%, when compared to the 3wt.% of a 319) allow this family of alloys to have excellent mechanical

conductivity.

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properties at room temperature, good damage tolerance and elevated heat

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Nowadays, T7 356 low copper primary alloys are not the main option anymore. T7 351 primary alloys with transition metals are able to reach higher

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mechanical properties (tensile and fatigue) than T7 356, from room temperature up to 300◦ C. Because of the elevated price of transition metal, only specific cylinder heads use 351 aluminum alloys. Moreover, higher mechanical properties are generally synonymous of lower elongation and then lower damage tolerance. Also, and because of extra alloying elements the heat conductivity of 351 alloys is typically lower than what 356 alloys offer. 3

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Alloy development has been the center of attention for the past ten years,

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particularly since computers can handle heavy computations. This has been emphasized by the use of Integrated Computational Materials Engineering (ICME) approach. Indeed, theoretical thermodynamic models are useful for

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cylinder heads at high temperature [1–8].

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developing new chemical compositions in order to improve the strength of

On the contrary, heat treatment development is an area that has not been thoroughly investigated. Optimized heat treatments could allow low copper 356 alloys to have better mechanical properties without losing damage tolerance properties and/or heat conductivity. During the last five years, the main

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development has been multi steps solution treatments. This allows the complete solutionizing of hardening elements, while limiting incipient melting and allowing extra globulisation of eutectic Si particles for better elongation and

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damage tolerance. Usually, multi steps solution treatments start at 495◦ C

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and go up to 540◦ C.

The main hardening systems in low copper 356 alloys are Mg2 Si and Q

phase (Al5 Cu2 Mg8 Si6 ) [9]. Precipitation starts with the nucleation of coherent Guinier-Preston zones (GP zones) in the pro-eutectic α-phase. When the size of the nucleus is larger than the critical radius of nucleation, GP zones will grow to eventually become GP2 zones [10–12]. Growth is driven 4

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by reduction of the supersaturation of the matrix and controlled by diffusion. Metastable GP2 zones eventually grow into semi-coherent precipitates [13–

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20]. Finally, when supersaturation in the matrix no longer exists, coarsening

takes place and precipitates become stable [19, 21–23]. Coarsening is driven

by the reduction of the interfacial energy of the precipitates and controlled by

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diffusion. It is worth noting that at elevated temperatures spontaneous nucle-

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ation of semi-coherent particles has also been mentioned [24].

The purpose of this work is to design a new heat treatment in order to improve the mechanical properties of low copper 356 primary alloys. This heat treatment combines a multi step solution treatment with a quench and a multi step overaging treatment. This new type of heat treatment takes advantage

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of the evolution of the solubility limit of the matrix as a function of temperature in order to allow successive precipitation sequences to happen during multistep overaging. This leads into an increased volume fraction of precipitates

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and eventually higher mechanical properties without jeopardizing the relief of

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residual stress.

After a short presentation of the material and the heat treatments (part 2),

Transmission Electron Microscopy (TEM) investigations are presented in part 3. Associated mechanical properties are discussed in part 4 of this article.

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2. Materials and methods

minum alloy. Its nominal composition is given in table 1. % Fe

% Si

% Cu

% Zn

% Mg

% Mn

% Ti

min % wt.

-

6.5

0.45

-

0.3

-

0.1

nominal % wt.

0.07

7

0.5

0.035

0.35

0.035

0.125

max % wt.

<0.15

7.5

0.55

<0.07

0.4

<0.07

% Ni

% Sn

% Sr

-

-

0.008

0.025

0.025

0.01

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Elements

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The material used in this study is a primary low copper 356 foundry alu-

0.15

<0.05

<0.05

0.012

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Table 1: Nominal composition of the primary low copper 356 aluminum alloy.

Ingots are melted into a 80kg crucible at a temperature of 720◦ C, and the chemical composition is adjusted. Liquid metal is then degassed using nitrogen. The treatment of the metal is done during degassing by adding Foseco Coveral flux. Density of the metal is measured by means of a resid-

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ual pressure test. AFNOR (NF EN 1706-2010) tensile samples are cast into a metallic permanent die (cf. figure 1) and liquid metal is filtered through a metallic filter inserted into the mould. The temperature of the die is monitored

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by two thermocouples and samples are cast when the die reaches a temperature of 150◦ C. This allows the secondary dendritic arm spacing (SDAS) to

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be close to 20µm.

After casting, four different heat treatments are made:

• T64: solutionizing at 525◦ C during 3h to 8h + air quench + aging at 165◦ C during 1h to 6h;

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Figure 1: AFNOR metallic die used to cast the samples.

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• T6: solutionizing at 525◦ C during 3h to 8h + air quench + aging at 180◦ C during 1h to 6h;

• T7: solutionizing at 525◦ C during 3h to 8h + air quench + aging at

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200◦ C during 1h to 6h;

• T7R: solutionizing at 500◦ C during 1h to 5h and 540◦ C during 1h to 5h

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+ water quench + aging at 200◦ C during 1h to 6h. This heat treatment

is considered to be the reference of this study.

After heat treating, the microstructure of the alloy is mainly composed of

a proeutectic α-phase, globular eutectic silicon and very few iron intermetallic compounds as represented in figure 2. Finally, after the various heat treat7

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ments all the alloys were artificially overaged (AOA) during 100h at 150◦ C.

Figure 2: Microstructure of T7 primary low copper 356 aluminum alloy.

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2.1. TEM insvestigations

Thin foils were conventionally made by final electrolytic polishing of a

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disk of 3 mm diameter, that was previously thinned to a thickness of 100

µm using regular mechanical polishing, into a solution of nitric acid (1/3) in methanol (2/3) at -30◦ C using a 12V tension. Three to four thin foils were at least made for each heat treatment. TEM observations were done with a JEOL 2000EX operating at 200kV. Scanning Transmission Electron Microscopy (STEM) and Energy-Dispersive Spectroscopy (EDS) capabilities of 8

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a 200kV operating Tecnai F20 were used for chemical mapping of precipi-

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tates.

On the one hand, low magnification (x20k and x25k) observations were made to display coarse precipitation of incoherent Q phase. On the other

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hand, the hardening coherent Q, β00 , θ00 phases and the semi-coherent θ0

phase have been observed at higher magnification, typically between x100k

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and x250k. These different phases were identified by the position of their diffraction spots with respect to the spots created by the aluminum matrix.

The statistical evaluation of the precipitate density was carried out by image analysis at the same magnification for each type of precipitate regardless

[001]Al zone axis.

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of the heat treatment, using the ImageJ software on images in dark field in

2.2. Mechanical testing

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Tensile testing was performed using an electro-mechanic Instron 3382 system coupled to a 100kN Instron 2716-002 load cell. Extensometer is an

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Instron AutoX 750 high resolution automatic device. This system to belongs to class 05, which means it has +/-0.5% uncertainty on the measurements. Data acquisition was set to 100Hz and Instron Bluehill software was used to postprocess the data. Finally, strain rate was 4mm.min−1 . For each heat

treatment, five tensile samples were utilized.

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3.1. Reference heat treatment (T7R)

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3. TEM observations

The precipitation of β00 (Mg2 Si) and Q phase (Al5 Cu2 Mg8 Si6 ) occurs during

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the T7R heat treatment and enhances the mechanical properties of the alloy. Figure 3 shows one TEM dark field (DF) micrograph of β00 (Fig. 3.a) and one

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TEM bright field (BF) micrograph of Q phase precipitates (Fig. 3.b). On this same micrograph beta precipitates can also be seen with low contrast. These two micrographs were taken after T7R heat treatment. It is worth noting that Al2 Cu particles do not seem to be present in the proeutectic α-phase in any of the four heat treatments, i.e. before overaging. This can be explained by

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the low amount of copper in the material and the presence of Q-phase.

3.2. Coarse precipitation

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Similar looking coarse and incoherent precipitates are present in all airquenched tempers (T64, T6 and T7 overaged during 100h at 150◦ C) but are

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absent in the water-quenched tempers (T7R and T7R overaged during 100h at 150◦ C). After a T6 air quench treatment, these precipitates have a size of 0.5 to 1 µm as it is displayed by the BF micrograph on figure 4.a. Dark

contrast bright-field STEM image (Fig. 4.b) and EDS elemental maping distribution show that these precipitates are mostly composed of Cu, Mg and Si (cf. fig. 4.c). Therefore, one can conclude that they belong to the Q phase 10

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(a)

(b)

Figure 3: TEM micrographs of a) DF of β00 (Mg2 Si) and b) BF of Q phase (Al5 Cu2 Mg8 Si6 ) after T7R on (002) zone axis (ZA).

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family.

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After overaging 100h at 150◦ C, the number of incoherent Q precipitates is lower for the T7 treatment (aged at 200◦ C) than for the T64 and T6 treatments (respectively aged at 165 and 180◦ C). Since the dissolution tempera-

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ture of the Q phase is lower than the solutionizing temperature [25], and these

coarse precipitates do not appear during aging for T7R specimens, this can

3.3. Fine precipitation 3.3.1. Q phase

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be attributed to the slower cooling rate of air quenching vs. the water quench.

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All the overaged tempers exhibit the precipitation of a coherent nanometric and spherical Q phase whose percentage and size remain almost constant and close to the reference temper after overaging, as it is shown in the dark field micrograph in figure 5.a. The different heat treatments and overag-

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ing treatments do not seem to influence the shape, the size or the percentage

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of the nanometric Q phase.

3.3.2. β00 phase

Unlike Q phase, β00 phase seems to be heat treatment and overaging de-

pendent. Indeed, on the overaged T64 temper, β00 is absent or in such a small number/size that it can not be distinguished using conventional TEM. 12

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(a)

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(b)

(c)

Figure 4: TEM micrographs of a) BF of coarse incoherent Q phase (Al5 Cu2 Mg8 Si6 ) for the T6 + overaging, (002) ZA b) STEM image of Q phase for the same HT and c) associated STEM-EDS Cu, Mg and Si elemental maps

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(a)

(b)

Figure 5: TEM micrographs of a) DF of coherent Q phase (Al5 Cu2 Mg8 Si6 ) for the T6 and overaging, (002) ZA b) DF of β00 (Mg2 Si) after T7R and overaging (2 variants).

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As a conclusion, the combination of aging time and temperature during the

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T64 heat treatment is not enough to achieve the formation of the β00 phase.

While increasing aging temperatures (180◦ C for T6 and 200◦ C for T7) the formation of β00 precipitates becomes more pronounced. However, it remains

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weak for the T6 when compared to both T7 tempers. For both T6 and T7

air quench temper the formation of β00 is lower than for the reference T7R

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water quench temper (cf. DF of figure 5.a). This can be explained by the highest cooling rate of the water quench that allows the matrix to be more supersaturated in solid solution. This, of course, has a direct impact on the formation of β00 precipitates after aging and overaging. It is worth noting that the diffraction spots of the Q and β00 phases are very close to the (200)Al zone

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axis on the diffraction pattern. Finally, the dark field micrograph presented in figure 5.b corresponds to the reference state (T7R) overaged 100h at 150◦ C.

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Two families of β00 precipitates oriented at 90◦ from each other can be seem.

3.3.3. θ0 phase

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After overaging 100h at 150◦ C, samples T64 and T6 show no θ0 precipi-

tates.

Heat treatments T7 and T7R have the highest aging temperature, i.e.

200◦ C. For both treatments, and after overaging 100h at 150◦ C, the formation of θ0 precipitates can be observed (cf. figure 6). Platelets shaped θ0 15

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precipitates are seen edge-on in the [200]Al diffraction vector or (002) zone axis on figure 6. After overaging 100h at 150◦ C, the size and number of θ0

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precipitates is slightly lower for the T7 air quenched temper than for the T7R water quenched temper as it is respectively shown in figures 6.a and 6.b.

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The bright field micrograph in figure 6.c, is made at high magnification for the overaged reference temper T7R. It shows the simultaneous presence of

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spherical Q precipitates (small black dots), of the three families of θ0 precipitates oriented at 90◦ from each other and displaying a size of about 100 nm in diameter, according to the orientation relationship (001)θ0 //{002}Al . Finally, two families of small β00 rod shaped precipitates with very weak contrast can

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also be seen in figure 6.c.

Table 2 displays a summary of the observations and the amount of the precipitates of each family, measured by image analysis, is expressed as

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a number of precipitates per square micrometer assuming the thickness of the thin foils is constant and the size of each family of precipitates is almost

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independent of heat treatment. • For both water quenched samples (T7R and overaged T7R), large Q precipitates were not directly observed. However, this phase is unambiguously present in all air quenched samples (T64, T6 and T7 heat treatments). • Nanoscale Q precipitates are observed for all heat treatments. They 16

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(a)

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(b)

(c) Figure 6: TEM micrographs of a) DF of θ0 (Al2 Cu) after T7 and overaging b) DF of θ0 (Al2 Cu) 00 0 after T7R and overaging and c) BF of the three 17phases (Q, β and θ ) after T7R and overaging on (002) ZA diffraction pattern.

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have the same shape, are of similar size and comparable in number.

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• The number of β00 precipitates increases with quench cooling rate and aging temperature. However, for the water quench, overaging during 100h at 150◦ C does not actually start the coarsening of β00 precipitates,

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their number per surface unit remaining substantially constant.

• The θ0 precipitates are only observed for the overaged T7 and T7R

T7R —

fine Q precipitates

1950± 180

Q size (nm2 )

4.1± 1.8

β00 θ0

T64 + AOA

T6 + AOA

T7 + AOA

T7R + AOA

+++

+++

++

—

2205±250

2410±150

1960±125

2395±480

3.9±1.7

4.2±1.7

4.4±2.0

4.6±1.8

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coarse Q phase

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conditions.

1875± 225

?

120± 23

105± 16

1882± 130

—

—

—

7.5± 3.2

24± 5.4

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Table 2: Summary of the TEM observations, the results are expressed in number of precipitates per square micrometer. T7R stands for T7 reference, and AOA for artificial overaged.

4. Mechanical properties and discussion Figure 7 displays the tensile properties at room temperature of the T64,

T6, T7 and T7R alloys before and after overaging 100h at 150◦ C.

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Before overaging, as one can expect the T6 heat treatment provides the best yield strength (YS=223.0±5.8MPa) amongst the air quenched samples.

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Then follows the T7 sample with a YS= 214.5±5.6MPa and finally the T64 sample with YS=210±5.5MPa. The ultimate tensile strength (UTS) provided by the three air quenching heat treatments are very similar: 287.6±5.2MPa

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for T64, 291.0±5.3MPa for T6 and 282.0±5.1MPa for T7. Due to data dispersion, the effect of heat treatment is less pronounced with respect to elonga-

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tion (EL) even though a trend can be observed: 7.0±0.5% for T64, 6.4±0.7% for T6 and 7.1±0.9% for T7. This is in good agreement with what could be expected from a T6 treatment when compared to a T64 or a T7 treatment. Finally, it is worth noting that the T64 and T7 heat treatments are providing

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almost identical elongation values.

The quality index QI [26] of the T64, T6 and T7 samples are respectively 415MPa, 412MPa and 409MPa. These values are very close to each other

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and show a good consistency in term of quality of casting and microstructure

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between the different samples.

Finally, as expected, the T7R sample displays higher mechanical proper-

ties because of the water quench: YS=258.0±6.8MPa, UTS=316.0±5.7MPa and EL=7.0±1.0%. This also gives a QI of 443MPa.

After overaging 100h at 150◦ C, the tendency is the same but the mechan19

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ical properties are higher:

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• for the T64 heat treatment: YS increases by 12% (236.0MPa vs. 210.0MPa), UTS increases by 6% (305.6MPa vs. 287.6MPa) and EL increases by 10% (7.7% vs. 7.0%);

YS increases by 6% (237.0MPa vs.

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• for the T6 heat treatment:

223.0MPa) and UTS increases by 3% (300.0MPa vs. 291.0MPa). How-

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ever, EL decreases by 22% (5.0% vs. 6.4%);

• for the T7 heat treatment: YS increases by 10% (236.4MPa vs. 214.5MPa) and UTS increases by 5% (295.0MPa vs. 282.0MPa). Similarly to above, EL decreases by 4% (6.8% vs. 7.1%);

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• for the T7R heat treatment: YS increases by 7% (275.6MPa vs. 258.0MPa), UTS increases by 4% (329.7MPa vs. 316.5MPa) and EL increases by 9% (7.6% vs. 7.0%);

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Figure 8 shows the evolution of the solubility limit of copper and magnesium in the aluminum matrix as a function of temperature for the nominal

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composition of the alloy at equilibrium: • at 165◦ C (T64 aging temperature), the solubility of Cu is 0.022 at.% and the solubility of Mg is 1.682x10−03 at.%;

• at 180◦ C (T6 aging temperature), the solubility of Cu is 0.029 at.% and the solubility of Mg is 2.569x10−03 at.%; 20

9 ±6 .7 .0

32 ±5 316 .7 .0

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±5 21 .6 4.5

±5 22 .8 3.0

±5 210 .5 .0

UTS (MPa)

±0 7. .9 1

±1 7 .0 .0

T7

T7R

±0 6.4 .7

T6

T64

T7R

T7

T6

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YS (MPa)

T64

T7R

T7

T6

T64

±0 7. .5 0

±0 5. .8 0

±0 6.8 .9

±0 7. .8 7

±1 7. .3 6

±5 30 .4 0.0

±5 29 .3 5.0

After overaging

±5 28 .1 2.0

±5 29 .3 1.0

±5 28 .2 7.6

±6 236 .2 .4

Before overaging

±6 25 .8 8.0

±6 237 .2 .0

±6 236 .2 .0

±7 275 .2 .6

±5 30 .5 5.6

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El (%)

Figure 7: Summary of the static mechanical properties measured before and after AOA for T64, T6, T7 and T7R heat treatments.

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• at 200◦ C (T7 aging temperature), the solubility of Cu is 0.043 at.% and the solubility of Mg is 4.335x10−03 at.%; • at 150◦ C (overaging temperature), the solubility of Cu is 0.015 at.% and

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the solubility of Mg is 1.068x10−03 at.%.

As a consequence, the amount of solid solution supersaturating the pro-

eutectic α-phase at the begining of overaging is: • for the T64 treatment: 0.0790 at.% in Cu and 0.0124 at.% in Mg; • for the T6 treatment: 0.0711 at.% in Cu and 0.0115 at.% in Mg; • for the T7 treatment: 0.0790 at.% in Cu and 0.0147 at.% in Mg. 21

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The increase in mechanical properties after overaging can mostly be explained by the secondary supersaturation of copper solid solution rather than

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magnesium. Indeed, after both T7 heat treatments and at the begining of the overaging at 150◦ C the theoretical secondary supersaturation in Cu is

0.0277 at.% whereas the theoretical secondary supersaturation in Mg is only

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3.27x10−03 at.%. As it can be expected, because of the lower aging temperatures for the T64 and T6 conditions, the secondary supersaturation in Cu

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and Mg is lower than for both T7 treatments. Even though 150◦ C is a quite low temperature, overaging during 100h allows a secondary precipitation of

θ0 particles to happen for all heat treatments. However, this secondary precipitation has only been seen for the T7 conditions using conventional TEM, as discussed in section 3. Nonetheless, the extra-hardening seen on the

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overaged T64 and T6 samples must come from the same phenomenon.

The secondary precipitation phenomenon has been modeled by means of

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a Kampmann-Wagner-Normalized (KWN) approach using homogeneous nucleation, growth and coarsening of θ precipitates on a 356+0.5wt.%Cu alloy.

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This model has been widely presented in [7]. Figure 9 shows the temperature input and the results of the KWN model, considering a single step and a double step aging. In order to be representative of what can be done at an industrial scale, including high volumes of production and cost savings, it was decided to model the following heat treatments (Fig 9.a). Indeed, the previous treatments, including a step at 150◦ C during 100h, are not industrially 22

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0.0020

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0.0016 0.0014 0.0012 0.0010 0.0008

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0.0006 0.0004 0.0002 0.0000

0

50

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at% of Cu in FCC-Al

0.0018

100 150 200 250 300 350 400 Temperature (°C) (a)

0.0020

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0.0016 0.0014 0.0012

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0.0010 0.0008 0.0006 0.0004

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at% of Mg in FCC-Al

0.0018

0.0002 0.0000

0

50

100 150 200 250 300 350 400 Temperature (°C) (b)

Figure 8: Equilibrium calculation of the solubility limit of a) copper ; b) magnesium in the aluminum matrix as a function of temperature for the nominal composition of the alloy. Com23 database. putation made with ThermoCalc using the TCAL4

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viable.

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• T7 water quench with single step aging: 4h at 200◦ C, which is comparable to the T7R temper;

• T7 water quench with dual step aging: 2h at 200◦ C followed by 2h at

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170◦ C.

As it can be seen on figure 9.b, the KWN model pictures the primary and

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secondary precipitations. During the quench, the supersaturation (in mol.% in figure 9.b) of the matrix increases and new Al2 Cu particles start nucleating. This is what can be called the primary supersaturation and precipitation. The particles larger than the critical radius of nucleation will then grow whereas the smaller particles will dissolve. The growth of the primary θ precipitates

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happens during the 200◦ C aging plateau and the primary supersaturation decreases. Soon after, the saturation of the matrix reaches the solubility limit at 200◦ C. This is when the coarsening of the primary precipitates happens.

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For the single step treatment, the volume fraction of the primary precipitates remains constant during coarsening until the end of the T7 treatment. How-

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ever, for the dual step aging T7 treatment when temperature goes from 200◦ C to 170◦ C, the supersaturation in copper of the pro-eutectic alpha phase increases again. This is the secondary supersaturation that occurs at the same time as the nucleation and growth of the secondary θ precipitates. Eventually, the saturation of the matrix reaches the solubility limit at 170◦ C and coarsen-

ing process takes over the growth. 24

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600

T7-single step aging Water quench

Maturation

Aging

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500

400

300

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temperature (°C)

T7-two steps aging

200

0 0.001

0.01

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100

0.1

1

t(h)

(a)

0.9

Volume fraction (%)

0.8 0.7 0.6 0.5

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0.4

Volume fraction-two steps aging

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Volume fraction-single step aging

0.1

0.01

EP

1.0

1

Oversaturation-two steps aging

0.001

supersaturation (mol.%)

Oversaturation-single step aging

0.3

0.2 0.001

0.01

0.1 t(h)

1

0.0001

(b) Figure 9: KWN model: a) input of temperatures ; b) evolution of the volume fraction of precipitates and supersaturation of the matrix as a function of time.

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These two T7 water quench heat treatments have been applied on samples, and figure 10 summarizes the measured mechanical properties. The

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secondary precipitation allows a gain of 6% in YS (269MPa vs. 255MPa) and 3% in UTS (332MPa vs. 322MPa). The decrease in elongation, which is

close to 33% (5.9% vs. 8.8%) is most likely the consequence of the harden-

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ing of the material combined to a statistical effect resulting from the foundry process and which can come from different sources such as the presence of

6. 5

2±

33

5 0. 8±

3 1. 9± 5.

T7 dual step

TE D

AC C

EP

T7 single step

8.

9± 26

25

5±

10

6. 4

32

2±

7. 5

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oxides, microshrinkage, more or less Sr in specific tensile specimens.

YS (MPa)

UTS (MPa)

El (%)

Figure 10: Summary of the static mechanical properties measured on samples after single and dual step aging.

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5. Conclusion

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The influence of a dual-step overaging is investigated on a primary 356+0.5wt%Cu aluminum alloy submitted to several heat treatments: T64

air quench, T6 air quench, T7 air quench and T7 water quench. After over-

aging 100h at 150◦ C, a secondary precipitation of Al2 Cu particles is seen

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using conventional TEM, mostly after the T7 treatments. This secondary pre-

cipitation leads to an increase of the static mechanical properties of the T7

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samples. After overaging 100h at 150◦ C, the T64 and T6 samples also display an increase of their mechanical properties but a secondary precipitation of Al2 Cu particles is not seen using conventional TEM.

The initial dual-step aging, with the second step of 100h at 150◦ C, has

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been adapted to an industrial context. Therefore, the following dual-step aging is considered: 2h at 200◦ C and 2h at 170◦ C. The latter has been compared to a more classic industrial T7 treatment with a single step aging set

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at 4h and 200◦ C. Both heat treatment are also modeled using a KWN model. The model captures the secondary oversaturation of the proeutectic α-phase

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that leads to the secondary precipitation during the dual-step aging treatment. This translates to an increase of the mechanical properties by 6% and 3% for the yield strength and the ultimate tensile strength respectively. The elongation decreases by 33%.

Dual step aging T7 heat-treatments can be useful when it comes to heat 27

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treating cylinder heads. On the one hand, it allows the static mechanical properties to be higher without impacting the relief of the residual srtresses

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during aging. On the other hand, this type of aging can be applied to any type of alloy and quench. A "pyramid" type heat treatment (i.e. increasing step

solution treatment + quench + decreasing dual-step aging treatment) could

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be a good solution for solving an issue which is typical to cylinder heads:

improving the damage tolerance of the part with a higher elongation (given

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by the step solution) and improving its resistance to crack initiation with a higher yield strength (given by the dual-step aging).

6. Acknowledgements

Thanks are due to Montupet S.A. for financial support and providing the

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samples.

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[1] M.J. Starink and A.-M. Zahra. An analysis method for nucleation and growth controlled reactions at constant heating rate. Thermochim. Acta, 292:159–168, 1997.

RI PT

[2] A. Deschamps and Y. Bréchet. Influence of predeformation and ageing of an Al-Zn-Mg

alloy - II. Modeling of precipitation kinetics and yield stress. Acta Mater., 47:293–305, 1999.

[3] O.R. Myhr and O. Grong. Modelling of non-isothermal transformations in alloys con-

SC

taining a particle distribution. Acta Mater., 48:1605–1615, 2000.

[4] O.R. Myhr, O. Grong, H.G. Fjaer, and C.D. Marioara. Modelling of the microstructure

Mater., 52:4997–5008, 2004.

M AN U

and strength evolution in Al-Mg-Si alloys during multistage thermal processing. Acta

[5] V. Vaithyanathan, C. Wolverton, and L.Q. Chen. Multiscale modeling of θ’ precipitation in Al-Cu binary alloys. Acta Mater., 52:2973–2987, 2004.

[6] R. Martinez, V. Russier, Couzinié, I. Guillot, and G. Cailletaud. Modeling of the influence of coarsening on viscoplastic behaviour of a 319 foundry aluminum alloy. Mater. Sci. Eng. A, 559:40–48, 2013.

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[7] R. Martinez, D. Larouche, G. Cailletaud, I. Guillot, and D. Massinon. Simulation of the concomitant process of nucleation-growth-coarsening of Al2 Cu particles in a 319 foundry aluminum alloy. Modelling Simul. Mater. Sci. Eng., 23(4):045012, 2015. [8] A. Assadiki, V.A. Esin, M. Bruno, and R. Martinez. Stabilizing effect of alloying elements

EP

on metastable phases in cast aluminum alloys by CALPHAD calculations. Computational Materials Science, 145:1–7, 2018.

AC C

[9] D. Ovono-Ovono. Recyclabilité des alliages d’aluminium de fonderie : influence des éléments résiduels sur la microstructure et le comportement mécanique. PhD thesis, Université de Technologie de Compiègne, France, 2004.

[10] W.A. Johnson and K.E. Mehl. Reaction kinetics in processes of nucleation and growth. Trans. AIME, 135:416–458, 1939.

[11] V. Gerold. Uber die Struktur der bei der Aushärtung einer Aluminium-Kupfer-Legierung auftretenden Zustande. Z. Metallkunde, 45:599, 1954.

29

ACCEPTED MANUSCRIPT

[12] X. Auvray, P. Georgopoulos, and J.B. Cohen. The structure of GP1 zones in Al-1,7% at. Cu. Acta Metall., 28:1061–1075, 1981.

1939.

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[13] M. Avrami. Kinetics of phase change I - General theory. J. Chem. Phys., 7:1103–1112,

[14] M. Avrami. Kinetics of phase change II - Transformation-time relations for random distribution of nuclei. J. Chem. Phys., 8:212–224, 1940.

ture. J. Chem. Phys., 9:177–184, 1941.

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[15] M. Avrami. Kinetics of phase change III - Granulation, phase change and microstruc-

[16] J.M. Silcock, T.J. Heal, and H.K. Hardy. Structural ageing characteristics of binary

M AN U

aluminium-copper alloys. J. Inst. Metals, 82:239–248, 1953-1954.

[17] G.W. Greenwood. The growth of dispersed precipitates in solution. Acta Metall., 3:243– 248, 1956.

[18] I.M. Lifshitz and V.V. Slyozov. The kinetics of precipitation from supersaturated solid solutions. J. Phys. Chem. Solids, 19:35–50, 1961.

[19] C. Laird and H.I. Aaronson. Mechanisms of formation of θ and dissolution of θ’ precipi-

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tates in an Al-4% Cu alloy. Acta Metall., 14:171–185, 1966. [20] G.C. Weatherly and R.B. Nicholson. An electron microscope investigation of the interfacial structure of semi-coherent precipitates. Philos. Mag., 17:801–831, 1968. [21] J.D. Boyd and R.B. Nicholson. The coarsening behaviour of θ00 and θ0 precipitates in

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two Al-Cu alloys. Acta Metall., 19:1379–1391, 1971. [22] D. Vaughan and J.M. Silcock. The orientation and shape of θ precipitates formed in

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Al-Cu alloy. Phys. Status Solidi, 20:725–736, 1967. [23] K. Hirano and S. Fujiwara. Impurity diffusion in aluminium. J. Nucl. Mater., 69-70:564– 66, 1978.

[24] J.M. Silcock and H.M. Flower. Comments on a comparison of early and recent work on the effect of trace additions of Cd, In or Sn on nucleation and growth of θ’ in al-cu

alloys. Scr. Mater., 46:389–394, 2002. [25] D. Ovono-Ovono, I. Guillot, and D. Massinon. The microstructure and precipitation

30

ACCEPTED MANUSCRIPT

kinetics of a cast aluminium alloy. Scr. Mater., 55:259–262, 2006. [26] M. Drouzy, S. Jacob, and M. Richard. Interpretation of tensile results by means of

AC C

EP

TE D

M AN U

SC

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quality index and probable yield strength. AFS Int. Cast Metals Jnl, 5:43, 1980.

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