Microstructure, tensile properties and wear resistance correlations on directionally solidified Al-Sn-(Cu; Si) alloys

Accepted Manuscript Microstructure, tensile properties and wear resistance correlations on directionally solidified Al-Sn-(Cu; Si) alloys Felipe Bertelli, Emmanuelle S. Freitas, Noé Cheung, Maria A. Arenas, Ana Conde, Juan de Damborenea, Amauri Garcia PII:

S0925-8388(16)33882-8

DOI:

10.1016/j.jallcom.2016.11.399

Reference:

JALCOM 39896

To appear in:

Journal of Alloys and Compounds

Received Date: 3 November 2016 Revised Date:

24 November 2016

Accepted Date: 29 November 2016

Please cite this article as: F. Bertelli, E.S. Freitas, N. Cheung, M.A. Arenas, A. Conde, J. de Damborenea, A. Garcia, Microstructure, tensile properties and wear resistance correlations on directionally solidified Al-Sn-(Cu; Si) alloys, Journal of Alloys and Compounds (2016), doi: 10.1016/ j.jallcom.2016.11.399. 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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Microstructure, tensile properties and wear resistance correlations on directionally solidified Al-Sn-(Cu; Si) alloys

Felipe Bertellia,b, Emmanuelle S. Freitasa; Noé Cheunga, Maria A. Arenasc, Ana Condec, Juan de

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Department of Manufacturing and Materials Engineering, University of Campinas, UNICAMP, 13083–860, Campinas, SP, Brazil

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Damboreneac, Amauri Garciaa

Department of Mechanical Engineering and Postgraduate Program of Mechanical Engineering, Santa Cecília University – UNISANTA, 11045-907, Santos, Brazil

National Center for Metallurgical Research (CENIM-CSIC), Avda Gregorio del Amo 8, 28040, Madrid, Spain

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Abstract

The development of alloys suitable for engine bearings demands not only reducing the amount of wear as well as increasing the load carrying capability due to both engines stop/start systems and sudden rise in load or velocity. Al-Sn based alloys are well-known for having excellent tribological and mechanical properties fulfilling these requirements: Sn is a self-lubricating component and addition of third elements should increase the strength of the Al-rich matrix. The

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current study focuses on interrelations of microstructure of directionally solidified Al-Sn(Cu;Si) alloys and mechanical/ tribological properties. In order to analyze the influence of alloy Sn content on the tribological behavior of these ternary alloys, ball-on-disc wear tests were performed under dry sliding conditions. Correlations between tensile strength,

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elongation and wear volume and the primary dendritic arm spacing (λ1) have been determined. The worn samples were analyzed by scanning electron microscopy and energy dispersive spectroscopy (SEM-EDS) and the wear scar

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topographies by a confocal profilometer. The analysis of the worn surfaces revealed a change from abrasive to adhesive wear mechanism for Al-Sn-Cu alloys and only adhesive one for Al-Sn-Si alloys. For Al-Sn-Cu alloys, the best wear resistance was observed to be related to coarser microstructures whereas refined microstructures improved the tensile properties, indicating an inverse trend between tensile and tribological properties. The wear resistance of the Al-Sn-Si alloys were shown not to be affected by the size of λ1, however the tensile strength is shown to increase significantly with the decrease in λ1. Keywords: Al-Sn Alloys, Directional solidification, Microstructure, Tensile properties, Wear.

_______________________________________________________________________ (*) Corresponding author E-mail address: [email protected]

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ACCEPTED MANUSCRIPT 1. Introduction

The main trend in automotive engineering is improving engine efficiency and fuel economy. This includes engine downsizing via lower weight requiring new materials for lightening structures

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with improved mechanical properties. Enhancing load-carrying capability needs bearing alloys with advanced tribological properties. The conventional white metal babbitt alloys do not hold up well for most applications in automobile engines due to their poor strength. Most bearing alloys consist of

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hard phases uniformly embedded in a soft matrix, soft phases in a hard matrix or a contiguous mixture of both. Such structure patterns have been found to possess the mechanical, friction and

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wear behavior that technically justify their bearing applications: the soft constituents behave as solid lubricants reducing friction at the tribocontact while the hard ones carry the load [1,2]. Traditionally, both Al-Pb and Al-Sn alloys have been used as self-lubricant bearing alloys, but, due to the environmental restrictions to the use of Pb in manufactured products, the search for alternative alloys

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has been concentrated on Al–Sn based alloys [3-5]. In selecting wear-resistant materials, it is of great importance to designers and engineers the understanding of the influence of the local microstructure on the wear resistance. Several experimental studies have investigated the dependence of wear

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mechanisms on microstructure features [6-12].

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Bearing performance of Al-Sn based alloys is strongly affected by the casting method as it affects the microstructures and properties of the alloy. Generally, during solidification, the Al-rich phase growths dendritically and Sn, due to its extremely low miscibility in Al, is contained within the interdendritic region forming discrete pools or pockets. Additions of Cu and Si to the Al-Sn alloys have potential to enhance the mechanical strength, wear resistance and the performance in seizure resistance [13-16]. Recent works have demonstrated that the refinement of microstructures is usually beneficial to mechanical properties in aluminum castings, justified by a homogeneous distribution of second/intermetallic phases acting as reinforcement components along the interdendritic and

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ACCEPTED MANUSCRIPT interphase areas [17-22]. Similarly, in terms of tribological behavior, it has been reported that a more homogeneous distribution of the self-lubricating particles of Al-Bi based alloys leads to lower wear volume, i.e. to optimum bearing properties [10, 23]. For the binary Al-Sn hypoeutectic alloys, an opposite trend has been observed as shown by the work of Cruz et al. [9]: when the dendritic

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network becomes coarser and consequently a higher amount of Sn is interdendritically segregated, the wear resistance was shown to improve. A similar behavior has been reported by Freitas et al. for Al-Pb and Al-In alloys, for which coarser droplets of the lubricating phase provided a more extensive

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protective film during the wear tests [10,11]. Therefore, with the increase of complexity caused by the addition of a third element to the Al-Sn system, a more thorough investigation needs to be

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performed on the perspective of a deeper understanding of the effects of microstructure features on the wear mechanisms.

In a previous study [24] a careful investigation on the microstructural evolution of Al-Sn(Cu;Si) ternary alloys, under a range of solidification cooling rates, was undertaken by the some of

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the present authors. These alloys were shown to be characterized by an Al-rich matrix of dendritic morphology, having the interdendritic regions formed by segregated Sn pockets and Al2Cu intermetallics or Si particles, respectively. The evolution of a representative length scale of the

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microstructure (primary dendritic arm spacing, λ1) along the growth direction of the directionally

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solidified (DS) alloys castings has been correlated with solidification thermal parameters, such as the growth rate (VL) and the cooling rate (TR), and experimental growth laws have been proposed, i.e.: Al-10wt.%Sn-10wt.%Cu: λ1 = 84VL-1.1; λ1 = 145TR-0.55 / Al-20wt.%Sn-10wt.%Cu: λ1 = 73VL-1.1 ; λ1 = 194TR-0.55; Al-15wt.%Sn-5wt.%Si: λ1 = 80VL-1.1 ; λ1 = 123TR-0.55 / Al-25wt.%Sn-5wt.%Si:λ1 = 125VL-1.0 ; λ1 = 204TR-0.55 , where λ1 [µ]; VL [mm/s] and TR [K/s].

The aim of the present investigation is to evaluate the effects of such parametric microstructural length scale (the primary dendritic arm spacing) of Al-Sn-Cu and Al-Sn-Si alloys on both tensile properties and tribological behavior under dry sliding conditions. Furthermore,

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ACCEPTED MANUSCRIPT experimental expressions correlating the ultimate tensile strength, yield tensile strength, elongation and wear volume with the primary dendritic arm spacing will be derived.

2.1. Directional solidification process and specimens for tensile tests

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2. Experimental

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The nominal compositions of the ternary alloys investigated in this work are: Al-

15wt.%Sn-5wt.%Si, Al-25wt.%Sn-5wt.%Si, Al-10wt.%Sn-10wt.%Cu and Al-20wt.%Sn-10wt.%Cu.

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The compositions of metals used to prepare these alloys are shown in Table 1. Solidification was carried out in a water-cooled apparatus, which promotes transient directional solidification. The casting assembly (upper part of Fig. 1: left side) and mold details have been described in the literature in previous studies by some of the present authors and collaborators [9,24]. The

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solidification setup permits a significant range of solidification microstructures to be obtained at different cooling rates in a single casting experiment. This is shown schematically in Fig. 1 (upper part/ right side), where higher cooling rates are associated with fine dendritic microstructures and

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vice-versa. As detailed in a previous study [24], the thermal data around the liquidus temperature of each alloy, provided by the cooling curves recorded by each thermocouple positioned along the

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length of the directionally solidified (DS) castings, were used to determine the coefficients of a 5thorder polynomial via the least square method. This has permitted temperature (T) vs. time (t) functions: T=f(t) to be generated. The derivative of these functions with respect to time yielded cooling rate functions TR =f(t). The experimental time corresponding to the liquidus front passing by each thermocouple is then inserted into the TR =f(t) function permitting the experimental cooling rate to be determined. The growth rate (VL) was based on the time-derivative of a fitting function representing a plot of position of each thermocouple against time of passage of the liquidus isotherm.

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ACCEPTED MANUSCRIPT The experimental values of VL and TR from top to bottom of the DS castings of the Al-Sn-Cu and Al-Sn-Si alloys, were shown to be in the range VL = 0.4 - 5mm/s and TR = 0.5 - 35K/s [24]. To analyze effects of microstructural features over mechanical and tribological properties of ternary Al-Sn-Cu and Al-Sn-Si alloys, samples were extracted from different positions along the

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length of the (DS) cylindrical castings. As shown in Fig. 1, (upper part- right side) the primary dendritic arm spacing, λ1 , increases from bottom to top of the DS casting. With a view to permitting a representative number of tensile and wear samples to be extracted, associated with equally

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representative λ1 mean values along the length of the DS castings, thin perpendicular samples have been extracted perpendicularly to the growth direction, as shown in the bottom part of Fig. 1.

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Physically, this means that the continuous evolution of λ1 along the length of the DS casting is being fitted by a step function, in which a particular sample is associated with each step. Thus, each sample will be associated with a different average value of λ1. The specimens for tensile tests were

2.2. Wear tests

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machined according to specifications of the ASTM Standard E 8M/04.

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The samples for wear tests were also extracted perpendicularly along the length of the DS

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castings, as shown in Fig. 1 (right side of bottom part) with a view to permitting a single mean primary dendritic arm spacing, to be associated with each sample, thus allowing the wear resistance to be correlated with the evolution of λ1. Unidirectional sliding wear tests were conducted using a with a ball-on-disk configuration. The tribological pair was formed by the Al-Sn(Si:Cu) alloys and a ceramic alumina ball (Al2O3) of 3 mm diameter. The normal load applied was 0.5N, which defined a mean contact pressure of ∼ 0.23GPa. Such value has been estimated considering the Young modulus and the Poisson ratio for the alumina of 300GPa and 0.21 respectively, and for the Al-Sn(Si:Cu) alloys the values are those described elsewhere [22]: 68GPa and 0.345, respectively. The diameter of

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ACCEPTED MANUSCRIPT the wear track was 1.5mm, the contact frequency 1Hz and the sliding distance 110m. Wear tests were performed under dry sliding friction conditions at a room temperature of 25 ±2 °C and relative humidity of 40%. The wear volume (V) was measured from the wear scar topographies obtained by an optical

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confocal profilometer Sensofar PLµ 2300, according to specifications of the ASTM G99-95a standard. Fig 2 shows schematically the ball-on-disc tribometer used in the wear tests and a sequence

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of images of post-process analysis. Both the tensile and wear tests were carried out in triplicate.

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3. Results and discussion

3.1. Microstructure features of Al-Sn-Si and Al-Sn-Cu alloys

To understand the microstructure formation during solidification of ternary Al-Sn-based

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alloys, two factors have to be taken into account: firstly, the solid solubility limit of Sn in Al is below 0.09 wt.%, which is extremely low; the second is the non-equilibrium kinetics of solidification that occurs under the transient heat flow conditions associated with the present directional solidification

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experiments. The microstructure of these alloys is characterized by a dendritic Al-rich matrix having

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tin pockets in the interdendritic regions. More details on the microstructural evolution of these alloys can be found in a previous study by some of the present authors [24], which correlates solidification thermal parameters and microstructure features. With the addition of a third alloying element, e.g. copper or silicon, monotectic/eutectic reactions give rise to the Al2Cu intermetallics or Si particles, respectively, which are also located in the interdendritic regions, as shown in optical (Fig. 3a/ Fig. 4a) and scanning electron microscopy (SEM) images (Fig. 3b/Fig. 4b) for Al-Sn-Si and Al-Sn-Cu alloys, respectively . These microstructures have also been reported in the literature for Al-Sn-Si

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3.2. Tensile properties vs. primary dendritic arm spacing

A number of studies reported in the literature highlighted the role of the length scale of cellular and dendritic microstructures on tensile properties of Al-based alloys. Refined

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microstructures having lower dendritic/cellular or interphase spacings, which are related to high growth/cooling rates during solidification, were shown to be associated with higher values of

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mechanical properties, such as hardness and tensile strength for different binary Al-alloys, e.g.: AlFe [ 25], Al-Sn and Al-Si [9] and ternary Al-Ag-Cu [17], Al-Zn-Mg [18], Al-Mg-Si [19], Al-Si-X [21] and Al-Cu-Co [26].

The addition of either copper or silicon into Al-Sn alloys is supposed to have a strengthening

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effect on the Al-rich matrix, which could be enhanced by an appropriate length scale of the dendritic microstructure. The tensile properties of Al-Sn-Si and Al-Sn-Cu alloys have been investigated as a function of a wide range of primary dendritic arm spacings. The evolution of λ1 with the position in

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the castings has been reported in a previous study by some of the present authors [24]. The tensile

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specimens were extracted along the length of the DS castings, associated with the different values of λ1, as described in the experimental section. Figs. 5 and 6 show the results obtained for ultimate tensile strength (σu), yield strength (σy: 0.2% proof stress) and elongation (δ) as a function of λ1, for the Al-Sn-Si and Al-Sn-Cu experimentally examined, respectively. The graphs are composed by average values of triplicate tests with the error bars representing the range of maximum/minimum measurements. Furthermore, for comparison purposes on the effects of adding Cu and Si to these alloys, experimental growth laws from the literature correlating λ1 and mechanical properties of Al-

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ACCEPTED MANUSCRIPT Sn binary alloys (compositions in the range 15 to 30wt.%Sn) [9], have also been inserted in Figs. 5 and 6. It can be seen in Fig. 5 that the increase from 15 to 25wt.% Sn in the Al-Sn-5wt.%Si alloys do not show any influence on σu, σy and δ with λ1, that is, single experimental laws are applicable to

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both alloys investigated. It is worth noting that the effect of the size of λ1 on σu and σy is significant. From the highest experimental λ1 value (lowest λ1-0.5 – coarser microstructure) to the lowest one (highest λ1-0.5 - finer microstructure), σu and σy experienced increase of about 50% (80 to 120 MPa)

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and 30% (73 to 95 MPa), as shown in Figs. 5a and 5b, respectively. In the same experimental range

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of λ1 the ductility, represented by the elongation to fracture, shows the same trend, i.e. increased from 2 to 5%. When these properties are compared with the corresponding values of binary Al-Sn alloys, the increase in the tensile strength of the ternary Al-Sn-Si alloys is remarkable (e.g. for the lowest experimental λ1 , σu increases from about 55MPa to 120 MPa: Fig. 5a). In contrast, the

about 13% (Fig. 5c).

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maximum elongation of the ternary Al-Sn-Si alloys is about 5%, while that of binary Al-Sn alloys is

Similar trends can also be observed in Fig. 6 for the evolutions of σu, σy and δ with λ1 for Al-

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Sn-10wt.%Cu alloys, i.e. with the decrease in λ1 both σu and σy increase (Figs. 6a and 6b, respectively). However, as shown in Fig. 6c, the effect of λ1 on δ is only noticeable for the Al-

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20wt.%Sn-10%wt.%Cu alloy. It can be seen that Cu has a remarkable role in the increase of the tensile strength when σu and σy of the Al-Sn-10wt.%Cu alloys are compared with the corresponding values of binary Al-Sn alloys ((Figs. 6a and 6b). While the addition of 5wt.%Si to Al-Sn alloys is shown to increase σu in about twice, the addition of 10wt.%Cu promotes increase in σu of up to three times along a similar range of primary dendritic arm spacings. The observed increase in tensile strength with the refinement of the dendritic array is related to the way the interdendritic phases are distributed throughout the alloy microstructure, as also observed in previous studies [9, 17, 19, 21]. Higher solidification cooling rates promote simultaneous 8

ACCEPTED MANUSCRIPT refinement of both the dendritic matrix and second phases. The better distribution of the reinforcement phases, such as Si and Al2Cu particles, for Al-Sn-Si and Al-Sn-Cu, respectively, will induce a more efficient blockage to the displacement of dislocations during the slip process. Therefore, it can be said that the tensile strength is governed not only by the amount of Si or Cu in

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the composition of Al-Sn-(Si;Cu) alloys, but also by the length scale of the Al-rich matrix.

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3.3. Wear and microstructure features

Figs. 7 and 8 show a sequence of SEM images of wear tracks (same magnification) of

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samples extracted from different positions along the length of the DS castings of Al-Sn-Si and AlSn-Cu alloys, respectively. Along with each image, the identification of sample position (P) from the cooled bottom of the casting, the corresponding solidification cooling rate (TR) and average primary dendritic arm spacing (λ1), can be found. A visual comparison, shows that the Al-Sn-Cu alloys

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samples (Fig. 8) experience a gradual decrease in the wear track for positions farther away from the cooled bottom of the castings (P=0). This is more evident at the wear track profiles shown in Fig. 9, in which lower width and depth of wear tracks can be observed, thus indicating higher wear

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resistance. That is, as λ1 increases with the increase in P, the tin pockets located in the interdendritic

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regions become thicker, and the wear track decreases. Conversely, for the Al-Sn-Si ternary alloys (Fig. 7) this is not visually evident and apparently, the wear track is essentially the same, indicating that both λ1 and tin distribution had little effect over the wear resistance under the experimental conditions investigated.

For the Al-10wt.%Sn-10wt.%Cu alloy, the friction coefficient (the resistance to relative motion between the two bodies in contact during the wear test) for all samples examined was found to be 0.48, however, with the increase in the Sn content of the alloy (Al-20wt.%Sn-10wt.%Cu alloy) it decreased slightly to 0.42. For Al-Sn-Si alloys, a same friction coefficient of about 0.52 was found 9

ACCEPTED MANUSCRIPT for all samples of both alloys. There is no direct relationship between the friction coefficient and the wear rate and it should certainly not be used as a criterion to explain the wear behavior of a material as there are other parameters (contact pressure, roughness, hardness, ductility, oxides formation, lubricant regime, etc), which affects wear. There is a simple approach saying the higher the

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coefficient of friction, the higher the wear. So, it should be used with caution and just as an indicator. However, as it has been pointed out by Kato [27], it is useful to the understanding of wear

mechanisms. In this sense, the solidification microstructures might drastically alter the surface

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roughness and other physicochemical properties of the surfaces in motion and then the response of the tribo-system.

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The total wear volume losses of the Al Sn (Si, Cu) alloys samples have been calculated from the topography of the wear tracks obtained by confocal profilometry, Fig. 9. As it can be seen there is a change in the wear profiles with different positions of Al-Sn-Cu alloys samples. However, for Al-Sn-Si alloys, the profiles are rather similar without noticeable effect on the microstructural

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features. It can also be seen that Al-Sn-Cu alloys with higher Sn content (20 wt.%Sn) had worn tracks of lower depths, indicating a better wear resistance. Figs.10 and 11 show the evolution of wear volume (V) as a function of the primary dendritic

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arm spacing, for Al-Sn-Si and Al-Sn-Cu alloys, respectively. Four samples, characterized by

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different λ1 values, were chosen for each alloy composition and subjected to wear tests under the conditions detailed in the experimental section, with a view to evaluating the role of microstructural effects on the wear behavior.

For the Al-Sn-Si alloys, Fig. 10, the results obtained are in agreement with the previous observations, which indicated that both λ1 and consequently the distribution of Sn-pockets in the interdendritic regions had little or no effect over the wear resistance. Furthermore, the increase in the alloy Sn content from 15wt.% to 25wt.% resulted in slightly higher wear volume, that is, the increase in the self-lubricating element has not favored the wear resistance. This result is in agreement with 10

ACCEPTED MANUSCRIPT the literature that pointed out that for Al-Sn bearings the highest scuff resistance is associated with a maximum of 20 wt.%Sn [28]. Conversely, for both Al-Sn-Cu alloys examined it can be seen that V decreases with the increase in λ1. Experimental points in Fig. 11 have been fitted by appropriate functions in order to

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permit experimental equations correlating V to the inverse of the square root of λ1 to be derived. It is worth noting that the dendritic coarser regions are associated with larger Sn-pockets in the

interdendritic regions, a microstructural arrangement that is shown to increase the wear resistance

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when compared with refined microstructures. This is in agreement with a previous study on binary

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Al-Sn alloys, in which higher values of λ1 were shown to be associated with improved wear resistance [9].

Fig. 12 (a) and (b) presents the wear resistance (WR=1/V) and σu as a function of λ1 for AlSn-Si and Al-Sn-Cu alloys, respectively. In this figure, the ultimate tensile strength and wear resistance vs. λ1 are combined in a same plot, thus permitting appropriate ranges of both WR and σu

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to be associated with the corresponding values of λ1. The wear resistance of Al-Sn-Si alloys is shown not to be affected by the size of λ1 (Fig. 12a), however, a refined microstructure (higher λ1-0.5) for the

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Al-15wt.%Sn-5wt%Si alloy is shown to combine highest tensile strength and wear resistance. Further increase in the Sn content of the alloy to 20wt.% brings no improvements. For Al-Sn-Cu

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alloys (Fig. 12b), the best wear resistance is observed to be related to coarser microstructures (lower λ1-0.5), whereas refined microstructures improve the tensile strength, indicating an inverse trend between tensile and tribological properties. The best balance between the two properties can be found for the Al-20wt.%Sn-10wt.%Cu associated with intermediate values of λ1 (50-60µm).

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3.4. Wear mechanisms

Fig. 13 shows SEM (SE: secondary electrons/BSE: backscattered electrons) micrographs with

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details of worn out surfaces of samples subjected to wear tests. Although the active wear aspect in Al-Sn-(Cu; Si) alloys has been mostly adhesive, it can be observed for the Al-Sn-Cu alloys that a transition from abrasion to adhesive wear occurs (Fig. 13a and 13b). However, for the Al-Sn-Si

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alloys all the wear surfaces examined exhibited only aspects of adhesive wear (Fig. 13c). The highlighted areas show typical characteristics of each wear mechanism found in the samples of the

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alloys examined.

Fig. 13a illustrates the two-body abrasive wear for the samples extracted from the initial positions of the DS Al-Sn-Cu alloys castings, region of higher solidification cooling rates. These samples have refined microstructures with small dendritic spacings and the hard Al2Cu intermetallic

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phase well distributed in the interdendritic areas. During sliding, the Al2Cu phase is fragmented. It generates particles that act as abrasive particles on the surface, which lead to the appearance of twobody abrasive wear, identified in area 1 of Fig. 13, characterized by parallel scratch marks in the

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sliding direction. For higher magnification of area 1, a SE image shows the characteristic grooves of

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abrasive wear and plow grooves, and a BSE image shows the regions containing the Sn-rich phase (in white). This Sn-rich phase usually act as a solid lubricant [29-31], however, since it is very well distributed in the interdendritic areas (refined microstructure) it is not able to act as an efficient lubricating agent. Nevertheless, the Sn-rich phase visibly minimizes the wear grooves of abrasive action generated by the Al2Cu particles. It is worth noting that on samples that undergone abrasive wear, the occurrence of Sn whiskers can be noted, as shown in Fig 13a. The growth of Sn whiskers has been reported in the literature for Al-Sn-Si alloys [29, 32], where the growth mechanism was shown to be based on a

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that is, with evidences of wear adhesive mechanisms indicated in the highlighted area 1. For coarser microstructures it can be observed that the lubricating action of the Sn-rich phase is more effective than the abrasive action induced by the Al2Cu phase. This behavior can be attributed to the formation

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of pockets with higher fraction of lubricating agent (Sn). During sliding, these Sn rich pockets are smeared out on the surface and protect it serving as solid lubricant. The Sn pockets form a film that

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minimizes wear because they act involving the abrasive nature of the particles and change the wear predominant mechanism from abrasive to adhesive (the smeared Sn phase is highlighted in the BSE image). The role of solid lubricants are similar to that of liquids: to separate surfaces and provide shear path within the interface [33].

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For Al-Sn-Si alloys samples (Fig. 13c) it can be observed that only adhesive wear occurred in all samples examined, represented by the highlighted area 1. The change in the length scale of the microstructure, caused by different solidification cooling rates, is shown not to affect the lubricating

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action of Sn in samples of both alloys compositions examined. In the highlighted areas, regions of

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adhesive wear with the solid lubricant "Sn" smeared out on the sample surface can be observed. The wear resistance is shown not to be affected by microstructural refining and the wear mechanism is always the same along the entire length of the directionally solidified castings, i.e., adhesive wear. In Figs. 13 (b) and (c), areas with cracks can be observed, which resulted from the delamination process of layers formed by transfer and adhesion of wear particles, which by the combined action of the solid lubricant/Sn and the sliding process are impregnated on the surface of the material. These layers, from certain thickness, give rise to the delamination process, as also reported by some researchers [12, 33-35].

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4. Conclusions

The following conclusions can be drawn from the present experimental investigation: Refined microstructures, i.e., regions with smaller primary dendritic arm spacings and having

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•

homogenous dispersion of Sn and Al2Cu/Si particles for the Al-Sn-Cu and Al-Sn-Si alloys,

respectively, were shown to be associated with higher tensile properties. The increase from 15 to

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25wt.% Sn content in the Al-Sn-5wt.%Si alloy and from 10 to 20wt.%Sn to Al-Sn-10wt.%Cu alloy was shown not to affect the trend of evolution of σu with λ1, that is, single experimental equations

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were shown to be applicable to both alloys of each system (Al-Sn-Si/Al-Sn-Cu) investigated: Al-10/20wt.%Sn-10wt.%Cu:

σu = 65 + 917 λ1-0.5

Al-15/25wt.%Sn-5wt.%Si:

σu = 38 + 675 λ1-0.5

where σu [MPa] and λ1 [µm].

While the addition of 5wt.%Si to Al-Sn alloys was shown to increase σu in about twice, the

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•

addition of 10wt.%Cu promoted increase in σu of up to three times along a similar range of λ1. •

The wear resistance, represented by the inverse of the wear volume, was shown to increase

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with the increase in λ1 for Al-Sn-Cu alloys, that is, regions with coarser α-Al dendritic phase and

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interdendritic regions formed by larger Sn pockets and Al2Cu particles were shown to be the best microstructural arrangement for wear resistance. Experimental equations relating the wear volume (V) to λ1 have been proposed: Al-10wt.%Sn-10wt.%Cu:

V = 1.23 λ1-0.5 - 0.11

Al-20wt.%Sn-10wt.%Cu:

V = 0.32 λ1-0.5 - 0.023

where V[mm3] and λ1 [µm]. The wear resistance of Al-Sn-Si alloys was shown not to be affected by the size of λ1.

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Acknowledgements The authors acknowledge the financial support provided by FAPESP- São Paulo Research Foundation, Brazil (grants 2012/16328-2; 2013/15478-3; 2013/23396-7; 2014/50502-5), CNPq - The

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Brazilian Research Council and CSIC-Spanish National Research Council (Project i-link0944).

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References

[1] H.M. Kim, T. S. Kim, C. Suryanarayana, B.S. Chun, Microstructure and wear characteristics of

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rapidly solidified Al–Pb–Cu alloys, Mater. Sci. Eng. A 287 (2000) 59–65.

[2] V. Bhattacharya, K. Chattopadhyay, Microstructure and wear behaviour of aluminium alloys containing embedded nanoscaled lead dispersoids, Acta Mater. 52 (2004) 2293–2304. [3] H. Torabian, J.P. Pathak, S.N. Tiwari, On wear characteristics of leaded aluminium-silicon

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alloys, Wear 177 (1994) 47-54.

[4] B.K. Prasad. Sliding wear behaviour of bronzes under varying material composition, microstructure and test conditions, Wear 257 (2004) 110–123.

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[5] X. Liu, M.Q. Zeng, Y. Ma, M. Zhu, Wear behavior of Al–Sn alloys with different distribution of

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Sn dispersoids manipulated by mechanical alloying and sintering, Wear 265 (2008) 1857–1863. [6] G.C. Yuan, X.M. Zhang, Y.X. Lou, Z.J. Li, Tribological characteristics of new series of Al-Sn-Si alloys, T. Nonferr. Metal. Soc. 13 (2003) 774-780. [7] T. Marrocco, L.C. Driver, S.J. Harris, D.G. McCartney, Microstructure and properties of thermally sprayed Al-Sn-based alloys for plain bearing applications, J. Therm. Spray Techn. 15 (2206) 634-639. [8] R. Schouwenaars, V.H.Jacobo, A. Ortiz, Microstructural aspects of wear in soft tribological alloys, Wear 263 (2007) 727-735.

15

ACCEPTED MANUSCRIPT [9] K. S. Cruz, E. S. Meza, F.A.P. Fernandes, J.M.V. Quaresma, L.C. Casteletti, A. Garcia. Dendritic arm spacing affecting mechanical properties and wear behavior of Al-Sn and Al-Si alloys directionally solidified under unsteady-state conditions, Metall. Mater. Trans. 41 (2010) 972-984. [10] E.S. Freitas, A.P. Silva, J.E. Spinelli, L.C. Casteletti, A. Garcia, Inter-relation of microstructural

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features and dry sliding wear behavior of monotectic Al-Bi and Al-Pb alloys, Tribol. Lett. 55 (2014) 111-120.

[11] E.S. Freitas, J.E. Spinelli, L.C. Casteletti, A. Garcia, Microstructure-wear behavior correlation

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on a directionally solidified Al-In monotectic alloy, Tribol. Int. 66 (2013) 182-186.

[12] T.A. Costa, M. Dias, E.S. Freitas, L.C. Casteletti, A. Garcia, The effect of microstructure length

767-776.

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scale on dry sliding wear behaviour of monotectic Al-Bi-Sn alloys, J. Alloys Compd. 689 (2016)

[13] T. Desaki, Y. Goto, S. Kamiya, Development of the aluminum alloy bearing with higher wear resistance, JSAE Rev. 21 (2000) 321-325.

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[14] C.J. Kong, P.D. Brown, S.J. Harris, D.G. McCartney, The Microstructures of a thermally sprayed and heat treated Al-20wt.%Sn-3wt.%Si alloy, Mater. Sci. Eng. A 403 (2005) 205-2014. [15] H.R. Kotadia, E. Doernberg, J.B. Patel, Z. Fan, R. Schmid-Fetzer, Solidification of Al-Sn-Cu

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based immiscible alloys under intense shearing, Metall. Mater. Trans. 40 (2009) 2202—2211.

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[16] H.R. Kotadia, A. Das, E. Doernberg, R. Schmid-Fetzer, A comparative study of ternary Al-SnCu immiscible alloys prepared by conventional casting and casting under high-intensity ultrasonic radiation, Mater. Chem. Phys. 131 (2011) 241-249. [17] R.N. Duarte, J.D. Faria, C. Brito, N.C. Veríssimo, N. Cheung, A. Garcia. Length scale of the dendritic microstructure a ecting tensile properties of Al–(Ag)–(Cu) alloys, Int. J. Mod. Phys. B 29 (2015) 155026-1 - 155026-9.

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ACCEPTED MANUSCRIPT [18] E. Acer, E. Çadirli, H. Erol, M. Gündüz, Effect of growth rate on the microstructure and microhardness in a directionally solidified Al-Zn-Mg alloy, Metall. Mater. Trans. A 47 (2016) 30403051. [19] C. Brito, T. Vida, E. Freitas, N. Cheung, J.E. Spinelli, A. Garcia, Cellular/dendritic arrays and

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intermetallic phases affecting corrosion and mechanical resistances of an Al–Mg–Si alloy, J. Alloys Compd. 673 (2016) 220–230.

[20] R.V. Reyes, R. Kakitani, T. A. Costa, J. E. Spinelli, N. Cheung, A. Garcia , Cooling thermal

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parameters, microstructural spacing and mechanical properties in a directionally solidified hypereutectic Al–Si alloy, Phil. Mag. Lett. 96 (2016) 228-237.

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[21] H. Kaya, A. Aker, Effect of alloying elements and growth rates on microstructure and mechanical properties in the directionally solidified Al–Si–X alloys, J. Alloys Compd. 694 (2017) 145–154.

[22] X. Liu, M.Q. Zeng, Y. Ma, M. Zhu, Promoting the high load-carrying capability of Al–

294–295 (2012) 387–394.

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20 wt%Sn bearing alloys through creating nanocomposite structure by mechanical alloying, Wear

[23] T.A. Costa, E.S. Freitas, M. Dias, C. Brito, N. Cheung, A. Garcia, Monotectic Al–Bi–Sn alloys

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directionally solidified: Effects of Bi content, growth rate and cooling rate on the microstructural

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evolution and hardness, J. Alloys Compd. 653 (2015) 243–254. [24] F. Bertelli, C. Brito, I. L. Ferreira, G. Reinhart, H. Nguyen-Thi, N. Mangelinck-Noël, N. Cheung, A. Garcia, Cooling thermal parameters, microstructure, segregation and hardness in directionally solidified Al-Sn-(Si;Cu) alloys, Mater. Design 72 (2015) 31-42. [25] P.R. Goulart, J.E. Spinelli, N. Cheung, A. Garcia, The effects of cell spacing and distribution of intermetallic fibers on the mechanical properties of hypoeutectic Al-Fe alloys, Mater. Chem. Phys. 119 (2010) 272- 278.

17

ACCEPTED MANUSCRIPT [26] E. Çadirli, H. Kaya, D. Rabiger, S. Eckert, M. Gündüz, Effect of rotating magnetic field on the microstructures and physical properties of Al-Cu-Co ternary eutectic alloy, J. Alloys Compd. 647 (2015) 471–480. [27] K. Kato, Wear in relation to friction: a review, Wear 241 (2000) 152-157.

mechanism, J. Kones Int. Combustion Engines 11 (2004) 393-398.

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[28] A. Krzymien, P. Krzymien, Properties of the bearing alloys used for slide bearings of crank

[29] Z.C. Lu, M.Q. Zeng, Y. Gao, J.Q. Xing, M. Zhu, Improving wear performance of dual-scale Al–

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Sn alloys by adding nano-Si-Sn: Effects of Sn nanophase lubrication and nano-Si polishing, Wear 338-339 (2015) 258–267.

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[30] D. W. Gebretsadik, J. Hardell, B. Prakash, Friction and wear characteristics of different Pb-free bearing materials in mixed and boundary lubrication regimes, Wear 340-341 (2015) 63–72. [31] S.C. Zhang, Q.L. Pan, J. Yan, X. Huamg, Effects of sliding velocity and normal load on tribological behavior of aged Al−Sn−Cu alloy, Trans. Nonferrous Met. Soc. China 26 (2016) 1809-

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1819.

[32] Z.C. Lu, Y. Gao, M.Q. Zeng, M. Zhu, Improving wear performance of dual-scale Al–Sn alloys:

225.

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the role of Mg addition in enhancing Sn distribution and tribolayer stability, Wear 309 (2014) 216–

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[33] D.K. Dwivedi, Adhesive wear behaviour of cast aluminium–silicon alloys: Overview, Mater. Des. 31 (2010) 251-253.

[34] E. Rabinowicz, Friction and wear of materials. New York: John Wiley and Sons, 1965, pp. 70– 118.

[35] A. D. Sarkar, Wear of metals. England: Pergamon Press, 1976, pp. 50–160.

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ACCEPTED MANUSCRIPT Figure Captions

Fig 1 – Schematic representation of the solidification apparatus, length scale of the microstructure (primary dendritic arm spacing) in accordance with position in the DS casting and relative positions

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of samples extracted for mechanical tests.

Fig 2 – Schematic ball-on-disc tribometer and post-process analysis of confocal profilometer. Fig 3 – Typical microstructure of ternary Al-Sn-Si alloys: (a) optical microstructure emphasizing the

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dendritic matrix and (b) SEM image highlighting the interdendritic region.

Fig 4 – Typical microstructure of ternary Al-Sn-Cu alloys: (a) optical microstructure emphasizing

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the dendritic matrix and (b) SEM image highlighting the interdendritic region. Fig 6 - Experimental expressions correlating mechanical properties and primary dendrite arm spacing (λ1) for Al-Sn-Cu alloys: (a) ultimate tensile strength, (b) yield strength and (c) elongation. R2 is the coefficient of determination.

DS Al-Sn-Si alloys castings.

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Fig. 7 – Wear scar SEM images of samples extracted from different positions along the length of the

Fig. 8 – Wear scar SEM images of samples extracted from different positions along the length of the

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DS Al-Sn-Cu alloys castings.

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Fig 9 – Wear scar profiles of ternary Al-Sn-Si and Al-Sn-Cu alloys samples. The legend indicates the positions of samples from the cooled bottom of the castings and the lateral figure shows a 3D representation indicating the path from where the wear profile was scanned by the confocal microscope profilometer.

Fig. 10 - Wear volume as a function of the primary dendritic arm spacing for Al-Sn-Si alloys. Fig. 11 - Wear volume as a function of the primary dendritic arm spacing for Al-Sn-Cu alloys. Fig. 12 - Wear resistance (1/V) and σu as a function of the primary dendritic arm spacing (λ1) for ternary (a) Al-Sn-Si and (b) Al-Sn-Cu alloys. 19

ACCEPTED MANUSCRIPT Fig. 13 - Morphologies of worn surfaces: (a) abrasive wear mechanisms of refined Al-Sn-Cu alloys samples; (b) adhesive wear mechanisms of coarse Al-Sn-Cu alloys samples, and (c) adhesive wear mechanisms of Al-Sn-Si alloys samples. SE: secondary electrons/BSE: backscattered electrons.

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Table Captions

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Table 1 – Chemical composition (wt.%) of metals used to prepare the alloys.

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Table 1 – Chemical composition (wt.%) of metals used to prepare the alloys Sn 0.005 balance -

Si 0.055 balance -

Cu 0.010 0.004 balance

Fe 0.073 0.008 0.320 0.014

Zn 0.05 -

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Ni 0.006 0.0001 0.010 -

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Al balance 0.005 0.110 -

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Element Al Sn Si Cu

Pb 0.006 0.047 -

Ca 0.002

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Fig. 3

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Fig. 4

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140

R2=0.9

120 110

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100

80 70

55

45 0.06

0.07

0.08

0.09

0.10

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Al-15wt%Sn-5wt%Si Al-25wt%Sn-5wt%Si σu= 38+675.λ−0.5 1 Al-Sn [9]

90

0.11

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Ultimate Tensile Strength(σu)

130

−0.5 1

λ

0.12

0.13

(µm)

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R2=0.82

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110 105 100 95 90 85 80 75 70 65

Al-15wt%Sn-5wt%Si Al-25wt%Sn-5wt%Si σy = 49+368.λ−0.5 1 Al-Sn [9]

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Yield Strength (σy) [MPa]

(a)

30 0.06

0.07

0.08

0.09

λ

−0.5 1

(b)

0.10

(µm)

0.11

0.12

0.13

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Al-15wt%Sn-5wt%Si Al-25wt%Sn-5wt%Si δ = -1.5+53.λ−0.5 1 Al-Sn [9]

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Elongation(δ)[%]

12

R2=0.72 6 4

0 0.06

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λ−0.5 (µm) 1

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240 220 2

R =0.9

160

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180

Al-10wt%Sn-10wt%Cu

140 120

Al-20wt%Sn-10wt%Cu −0.5 σu= 65+917.λ1

100

Al-Sn [9]

60

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200

55

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0.12

0.13

−0.5

λ1

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0.15

0.16

(µm)

(a)

150

120 110 100 90

2

R =0.81

2

R =0.96

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Al-10wt%Sn-10wt%Cu −0.5 σy = 96+257.λ1

40 38

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130

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Al-20wt%Sn-10wt%Cu −0.5 σy = 63+410.λ1 Al-Sn [9]

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34 0.08

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δ=1

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Al-20wt%Sn-10wt%Cu δ = -1.9+32.λ−0.5 1 Al-Sn [9]

12 11

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Elongation(δ)[%]

Al-10wt%Sn-10wt%Cu 14

10

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Al-15wt.%Sn-5wt.%Si P = 10mm

P = 30mm

TR = 7.5K/s

TR = 3.4K/s

TR = 1K/s

λ1 = 41 µm

λ1 = 62 µm

λ1 = 123 µm

P = 50mm

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P = 5mm

TR = 0.55K/s

Al-25wt.%Sn-5wt.%Si P = 10mm

P = 30mm

P = 50mm

TR = 16K/s

TR = 7.4K/s

TR = 2.2K/s

TR = 1.3K/s

λ1 = 45 µm

λ1 = 68 µm

λ1 = 130 µm

λ1 = 178 µm

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P = 5mm

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λ1 = 170 µm

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Al-10wt.%Sn-10wt.%Cu P = 30mm

TR = 18K/s

TR = 10K/s

TR = 3.6K/s

λ1 = 29 µm

λ1 = 41 µm

λ1 = 71 µm

TR = 2.3K/s

λ1 = 92 µm

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P = 5mm

P = 10mm

P = 30mm

P = 50mm

TR = 48K/s

TR = 21K/s

TR = 5.5K/s

TR = 3K/s

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Al-20wt.%Sn-10wt.%Cu

P = 50mm

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P = 10mm

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P = 5mm

λ1 = 75 µm

λ1 = 106 µm

λ1 = 36 µm

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λ1 = 23 µm

Fig. 8

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Al-15wt%Sn-5wt%Si

Al-25wt%Sn-5wt%Si

10

10

5

5

0 -5

-10

50mm 30mm 10mm 5mm

-15 -20

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Z(µm)

Z(µm)

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Al-10wt%Sn-10wt%Cu

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Tracks profile of wear scar (µm)

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Tracks profile of wear scar (µm)

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Al-15wt%Sn-5wt%Si Al-25wt%Sn-5wt%Si

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0.20

Al-10wt%Sn-10wt%Cu V = 1.23.λ−0,5 - 0.11 1

Al-20wt%Sn-10wt%Cu V = 0.32.λ−0,5 -0.023 1

R2=0.99

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0.00 0.08

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0.22

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Al-25wt% Sn-5wt% Si (1/V) Al-15wt% Sn-5wt% Si (1/V)

120

45 40

105

35 30

90

25 20

75

0.07

0.08

0.09

λ−0.5 (µm) 1

100

60 0.12

260

90

σu

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Al-10wt% Sn-10wt% Cu (1/V) Al-20wt% Sn-10wt% Cu (1/V)

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Ultimate Tensile Strength (σu)

10 0.06

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Ultimate Tensile Strength (σu)

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σu

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Al-Sn-Si;Cu alloys microstructure: dendritic Al-matrix, Sn pockets and Si; Al2Cu particles The lower the dendritic spacing, λ1, > the tensile strength of Al-Sn-(Cu;Si) alloys

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The higher λ1, the higher the wear resistance for Al-Sn-Cu alloys

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Experimental equations relating wear volume/tensile properties vs. λ1 are proposed

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•