Influence of preform preheating on dry sliding wear behavior of 304 stainless steel reinforced A356 aluminum matrix composite produced by melt infiltration casting

Accepted Manuscript Influence of preform preheating on dry sliding wear behavior of 304 stainless steel reinforced A356 aluminum matrix composite produced by melt infiltration casting Ridvan Gecu, Ş. Hakan Atapek, Ahmet Karaaslan PII:

S0301-679X(17)30328-6

DOI:

10.1016/j.triboint.2017.06.040

Reference:

JTRI 4800

To appear in:

Tribology International

Received Date: 3 May 2017 Revised Date:

24 June 2017

Accepted Date: 26 June 2017

Please cite this article as: Gecu R, Atapek ŞHakan, Karaaslan A, Influence of preform preheating on dry sliding wear behavior of 304 stainless steel reinforced A356 aluminum matrix composite produced by melt infiltration casting, Tribology International (2017), doi: 10.1016/j.triboint.2017.06.040. 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.

ACCEPTED MANUSCRIPT Influence of Preform Preheating on Dry Sliding Wear Behavior of 304 Stainless Steel Reinforced A356 Aluminum Matrix Composite Produced by Melt Infiltration Casting Ridvan Gecu1*, Ş. Hakan Atapek2, Ahmet Karaaslan1 *[email protected] 1

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Yildiz Technical University, Department of Metallurgical and Materials Engineering, Istanbul, Turkey 2

Kocaeli University, Department of Metallurgical and Materials Engineering, Kocaeli, Turkey

ABSTRACT

The effect of preform preheating on wear behavior of 304 reinforced A356 matrix composite was

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studied. Stainless steel shavings were pressed to attain porous preforms which were infiltrated by molten A356 alloy in a subsequent process. All casting operations were performed at 730˚C.

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Preforms were inserted into mould before casting and held for varied times from 0 to 60 min. Manufactured composites were characterized by optical microscope, SEM, FEG-SEM, XRD, EDS and ball-on-disc type tribometer. Double-layered intermetallic compounds consisting of Fe4Al13 and Fe2Al5 phases were formed at interfaces, except 60 min preheated specimen. Both inadequate and excessive preform preheating adversely affected on dry sliding wear resistance while sufficient

1. INTRODUCTION

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preheating led to improve tribological properties of 304 reinforced A356 matrix composite.

Metal matrix composites (MMCs) have been largely used in automotive, aviation, transportation and defense industry applications where high mechanical strength, high wear resistance, good thermal

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stability, good corrosion resistance and low density features are expected from the material [1]. With increasing demand for lightweight structures, ceramic reinforced MMCs come to the forefront owing to their high strength/density ratio. Despite providing beneficial properties such as high hardness,

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high specific strength and good wear resistance, ceramic reinforced MMCs cause decrease in ductility and toughness, in addition to undesirable phase transformations at the interface and wetting problems between reinforcement and matrix alloy [2]. On the other hand, metal reinforced MMCs have great potential to replace ceramic reinforced MMCs due to low cost environmentally friendly production without losing competitive properties. The first requirement of MMCs in most applications is low density that makes aluminum alloys widespread as matrix materials. There have been numerous studies on SiC [3], Al2O3 [4], CNTs [5], TiC [6], TiB2 [7], B4C [8], ZrB2 [9] and graphite [10] reinforced Al matrix composites while metal reinforcement use in Al matrix remains limited. Especially steel based metal reinforcements show themselves up to enhance mechanical and tribological properties of Al matrix. Baron et al. [11]

ACCEPTED MANUSCRIPT researched Al matrix composite with both unalloyed steel (UAS) and stainless steel (SS) reinforcements. SS reinforcement increased tensile strength considerably thanks to the interfacial phase formation which caused decrease in tensile strength at UAS reinforced composite. For SS reinforced sample, reaction phase connected SS and Al at the interface strongly, while same reaction phase acted brittle and decreased tensile strength in UAS reinforced one because of its excessive

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volume fraction. It can be concluded that there is an optimum value for reaction phase ratio at the interface. Good wetting is also crucial condition for the powerful bonding to allow load transfer from matrix to reinforcement without failure [12].

Al matrix composites can be produced by various techniques like powder metallurgy [13], spray

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deposition [14], squeeze casting [15] and stir casting [16]. The most inexpensive near-net shape MMCs are manufactured via molten metal routes. Melt infiltration casting is a molten metal process that can be used to produce composites containing high reinforcement volume fraction (>50%) with

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using appropriate porous preforms [17]. Moreover, melt infiltration into preform is possible even under pressureless conditions through vacuum assisted casting procedure [18]. It is well-known fact that ceramic reinforced Al matrix composites display better wear resistance than that of unreinforced Al alloys [19–21]. However, tribological behavior of metal reinforced Al matrix composites have not been studied as much as ceramic reinforced ones and much effort is required to

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obtain wear characteristics in detail. It should be noted that different types of reinforcements and different wear conditions may cause different wear mechanism [22]. Reinforcement affects wear resistance either positively or negatively depending upon its type, particle size, distribution and volume fraction. Moreover, with the common effects of corrosion and wear, weight losses in

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material can be increased dramatically [23]. For this reason, corrosive properties should be taken into account in material selection. 304 SS alloy was chosen as a reinforcement in this study owing to

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its excellent corrosion resistance which is mainly attributed to its high Cr content [24]. In this work, dry sliding wear behavior of 304 SS reinforced A356 Al matrix composite fabricated by melt infiltration casting was studied. SS shavings were used to obtain porous preform and molten Al was infiltrated into its vacancies under vacuum atmosphere. Ball-on disc type tribometer was conducted using Al2O3 ball and the effects of preform preheating on wear resistance were investigated. 2. EXPERIMENTAL PROCEDURE Melt infiltration was carried out into two steps consisting of mould making and casting. Molten wax was cast into cylindrical plastic mould at 85°C and wax pattern with the dimensions of 25 mm in diameter and 50 mm in height was obtained. Castable plaster slurry was prepared by mixing

ACCEPTED MANUSCRIPT investment powder with water at the water/powder ratio of 40%. Wax pattern was placed into stainless steel perforated flask and castable slurry was poured into flask under vibration. Subsequent to waiting under vibration for 15 min to set, the mould was held for 2 h in undisturbed condition to solidify. Dewaxing was performed at 110°C for 1 hour to obtain mould cavity and dewaxed mould

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was heated up to 700°C gradually. Heating regime of plaster mould was given in Fig. 1.

Figure 1. Heating regime of plaster mould

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After mould making step, casting operation was performed. A356 aluminum alloy and AISI 304 SS shavings were chosen as matrix and reinforcing materials, respectively. Chemical compositions of these alloys were reported in Table 1. A356 alloy was selected due to its remarkable fluidity and castability properties while 304 SS shavings were used to improve mechanical and tribological

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features of Al matrix. Nearly same sized SS shavings obtained from turning machine were pressed under 220 MPa in order to get porous one-piece preform with the dimensions of 20 mm in diameter

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and 10 mm in height. Produced preforms were settled into heated mould before casting for varied times from 0 to 60 minutes at the last step of mould heating regime. Preform preheating process led to avoid rapid solidification of molten Al alloy and interaction time between liquid Al and solid SS preform was increased.

Table 1. Chemical compositions of experimented alloys (wt.%) Alloy

C

Si

Mn

Cr

Mg

Cu

Zn

Ni

Ti

Fe

Al

A356

-

7.0

0.1

-

0.35

0.2

0.1

-

0.2

0.2

Bal.

304 SS

0.08

1.0

2.0

18.0

-

-

-

8.0

-

Bal.

-

Plaster mould with preform inside was taken out from furnace at 700°C and placed into the vacuum chamber as given in Fig. 2 which shows schematic illustration of melt infiltration casting. All casting

ACCEPTED MANUSCRIPT processes were carried out at 730°C under -105 Pa vacuum pressure for 20 min. SS preform cavities were infiltrated by molten A356 Al during this time. When solidification was completed, mould was water cooled and composite sample was taken out. Obtained composites contain A356 and 304 SS alloys at the ratio of 50% by volume (the weight percentage of 304 SS reinforcement ratio is

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approximately %75).

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Figure 2. Schematic illustration of melt infiltration casting For cross-sectional analysis of manufactured composites, pieces were cut out from the specimens and embedded in a cold resin. After grinding and polishing steps, cross-sectional micrographs were

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taken using light optical microscope (Nikon Eclipse MA100), SEM (JEOL JSM 6060) and FEG-SEM (JEOL JSM 6335F). Formed phases were identified by XRD (Philips PW 3710) using CuKα radiation over 2θ range of 10-90°. The Joint Committee on Powder Diffraction Standards (JCPDS) database was used

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for the classification of the obtained reflexes. EDS analysis was performed to attain elemental content in specific spots and lines. Formation and distribution of main elements in the microstructure were identified by mapping analysis via FEG-SEM (JEOL JSM 6335F). Ball-on-disc type tribometer was conducted on samples after microstructural investigations. Al2O3 ball with 6 mm diameter was used as counterpart. Load, sliding distance and rotating speed were selected as 10 N, 100 m and 0.1 m/s, respectively. Sliding distance dependent friction coefficient values were recorded by tribometer during wear test. For each sample, 3 wear tests under same conditions were performed in accordance with ASTM G133-05 and the average values of friction coefficients and wear rates were given. WR=V/(L.P) formula where WR is wear rate, V is volume loss, L is sliding distance and P is applied load was used to determine wear rates of composites. Volume losses were

ACCEPTED MANUSCRIPT computed from average cross-sectional area of the wear tracks and length of the strokes. The heights and the widths of wear tracks were measured by profilometer (MAHR S2) and image analyser assisted optical microscopy (Zeiss Axiotech), respectively. In order to compare and verify calculated wear rate results, mass losses were recorded after wear tests and volume losses were also measured by dividing these mass losses by composite density according to ASTM G99-17. All worn surfaces

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were examined by SEM (JEOL JSM 6060) and discussed as a function of microstructural features.

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3. RESULTS AND DISCUSSION

Figure 3. Micrographs of 304 SS reinforced A356 matrix composites with (a) 0 min, (b) 10 min, (c) 30 min and (d) 60 min preform preheating

The micrographs of produced 304 SS reinforced A356 alloy matrix composites with varied preform preheating times were given in Fig. 3. Double-layered intermetallic compounds named as IMC I and IMC II were obtained in 0, 10 and 30 min preheated samples (Fig. 3a, 3b and 3c) while only monolayer reaction phase (IMC II) was formed in 60 min preheated specimen (Fig. 3d). Excessive oxidation exposure during preheating caused rapid CrO2 film growth on SS surface of 60 min preheated sample. CrO2 film may act as a barrier to diffusion during infiltration and consequently weakened interaction between matrix and reinforcement may be responsible for the nonformation

ACCEPTED MANUSCRIPT of IMC I at the interface. Besides, the interfacial reaction phase uniformity of 60 min preheated specimen was less than that of other specimens. Its continuous layer was locally disrupted by casting cavities and divided into more than one layer. Similarly, casting cavities of nonpreheated sample were also distinguished at the interface. However, those cavities did not affect uniformity or

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continuity of reaction phases as much as they did in 60 min preheated one.

Figure 4. Low-magnification SEM images of 304 SS reinforced A356 matrix composites with (a) 0 min, (b) 10 min, (c) 30 min and (d) 60 min preform preheating SEM micrographs of manufactured composites with different preform preheating times were shown in Fig. 4 (lower magnification) and Fig. 5 (higher magnification). Contrary to optical microscope images, dark regions refer A356 Al alloy whereas light regions point out 304 SS reinforcement.

ACCEPTED MANUSCRIPT Overall morphology of the composites were clearly visible in Fig. 4. It can be said that good mechanical and metallurgical bonding were obtained at all interfaces, provided that a few minor casting defects were neglected. Continuous double-layered interfacial reaction phases were obviously seen in 0, 10 and 30 min preheated specimens (Fig. 5a, 5b and 5c). No phase formation on steel side was observed while nonuniform discontinuous phase formation on aluminum side was

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obtained in 60 min preheated sample (Fig. 5d). Also, aforementioned CrO2 film was clearly visible in this specimen. It covered SS surface and blocked Al infiltration into preform cavities. It can be inferred that CrO2 film would almost prevent formation of IMC II which was occurred only where no

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CrO2 was present.

Figure 5. High-magnification SEM images of 304 SS reinforced A356 matrix composites with (a) 0 min, (b) 10 min, (c) 30 min and (d) 60 min preform preheating

ACCEPTED MANUSCRIPT Phase identification of manufactured composites was determined using XRD with CuKα radiation. Xray diffractograms were taken from the cross-sections of specimen that preheated for 10 min. XRD analysis results were given in Fig. 6. Al, Fe, Fe4Al13, Fe2Al5 and CrO2 phases were obtained. CrO2 phase was naturally formed on all sample surfaces at room temperature after composite production. It should not be confused with 60 min preheated sample which already contained CrO2 at the SS

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surface while casting process had not been started yet. Experimentally calculated Fe-Al phase diagram [25] was demonstrated in Fig. 7. Al-rich portion of the diagram indicates that when molten aluminum contact with the iron for enough time, the first reaction phase is θ-Fe4Al13. There is a dilemma to determine stoichiometric composition of θ phase. Kattner and Burton [26] proposed the chemical formula of θ as FeAl3. In other words, they reported that θ phase consists of 75% Al and

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25% Fe. However, according to the study of Grin et al. [27] , Fe4Al13 was confirmed as a crystal structure of θ, instead of FeAl3. Han et al. [25] confirmed Grin’s report with their findings which

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indicate the Al amount of θ phase as 76.5%. As a result of these previous studies, θ was called as Fe4Al13 in this work. According to Fig. 7, when Fe content in Al increases, η-Fe2Al5 is the second intermetallic compound to form. η-Fe2Al5 phase formation depends on the peritectic reaction of θFe4Al13 → Liquid + η-Fe2Al5 . Considering micrographs in Fig. 3d and 5d which showed monolayer interfacial phase, it can be readily predicted that CrO2 formation prohibited peritectic reaction much

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more than θ formation.

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Figure 6. XRD results of 304 SS reinforced A356 matrix composite with 10 min preform preheating

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Figure 7. Experimentally calculated phase diagram of the Fe-Al binary system [25]

ACCEPTED MANUSCRIPT Figure 8. SEM image and EDS spots of 10 min preheated sample Table 2. Elemental content in wt.% of EDS spots in Fig. 8 Al

Fe

Si

Cr

C

1

56.44

28.26

-

9.454

5.836

2

59.84

18.62

9.395

5.959

6.178

3

56.15

29.444

-

8.597

5.809

4

59.59

18.17

9.985

5.978

6.270

5

56.32

28.39

-

8.189

7.088

6

59.225

20.185

9.579

5.270

5.740

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Spot

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After phase identification by XRD diffractograms, EDS analysis was conducted on samples to specify IMC I and IMC II phases. Fig. 8 shows EDS spots of 10 min preheated sample and Table 2 represents

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related EDS results. 3 spots were measured for each layer to ensure reliability of results. Spot 1, 3 and 5 involves less Al and more Fe than that of spot 2, 4 and 6. Considering phase diagram and EDS results together, it can be surely said that IMC I refers η-Fe2Al5 phase while IMC II implies θ-Fe4Al13. Layer in Al side (θ) contains Si in substantial amounts. Complex Fe-Si-Al intermetallic compound like FeSiAl3 may contingently be formed although XRD findings did not support this prediction. Fe2Al5 is a Cr-rich and Si-poor phase because of the proximity to SS side. Si shows a tendency to accumulate into

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Al compound formation.

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θ-Fe4Al13 instead of diffuse in η-Fe2Al5 layer. This tendency increases the possibility of complex Fe-Si-

ACCEPTED MANUSCRIPT Figure 9. Results from line scanning analysis in Fig. 5b Fig. 9 shows the results from line scanning analysis in Fig. 5b. Analysis direction was from up to the bottom. The portion of the steel side included in the measurement was approximately 20 µm. Fe and Cr amounts were naturally higher than Al and Si in there. When steel side line was finished and ηFe2Al5 line was begun, Si intensities showed no differentiation while Al amount was suddenly

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increased. In θ-Fe4Al13 line that roughly represents between 30 and 40 µm, Si amounts were increased dramatically. There was also slight increase in Al values while Fe and Cr intensities continued to decline. In Al side, both Al and Si values had peaks and troughs because A356 microstructure intrinsically consists of α-Al and eutectic (α-Al + Si) phases.

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SEM image and related line scanning analysis of 60 min preheated sample were given in Fig. 10. Line scan direction was from left to the right. This measurement only considers chromium and oxygen contents. The darkest layer as shown in Fig. 10a was CrO2 according to line scan data in Fig. 10b. Cr

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and O values were only peaked when this layer was subjected to analysis. It should also be noted that θ-Fe4Al13 phase was occurred in a significant amount where CrO2 layer did not exist. Moreover, it was formed on a limited scale even at the presence of CrO2. The left corner of Fig. 10a was an evidence of previous prediction of this work which argued that CrO2 has more inhibiting effect on η-

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Fe2Al5 phase comparing to θ-Fe4Al13 phase.

Figure 10. (a) SEM image and (b) line scanning results of 60 min preheated specimen

30 and 60 min preheated specimens were also examined by FEG-SEM for EDS mapping analysis. FEGSEM images and mapping analysis results of these specimens were given in Fig. 11. EDS mapping results of 30 min preheated sample showed that the distribution of Al and Fe was homogenous along reaction phases. Si atoms preferred to accumulate into interfacial region rather than Al matrix. Significant amount of Cr was diffused in θ-Fe4Al13 and η-Fe2Al5 layers from SS side. It can be deduced

ACCEPTED MANUSCRIPT from considering Fe and Cr maps together that almost every point where Fe is found also contains Cr. This was due to the presence of Cr in Fe lattice as a substitutional atom. For 60 min preheated specimen, oxide layer was observed at the interface instead of θ and η phases. According to EDS mapping data, oxygen amount was peaked at the darkest layer of Fig. 11b. This layer also includes Cr as a high level due to CrO2 formation. However, Cr content was dramatically decreased just above

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the layer which restricted to Cr diffusion beyond Al matrix. Si accumulation was obviously noticed just above the oxide layer. Si atoms preferring interface rather than Al matrix tended to accumulate in a region where they could not join a phase. Considerable amount of Ni was also obtained in overall

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microstructures of both 30 and 60 min preheated samples.

Figure 11. FEG-SEM images and EDS mapping analysis results of (a) 30 min and (b) 60 min preheated specimens

ACCEPTED MANUSCRIPT Layer thicknesses were also calculated via image analyser assisted light optical microscopy to determine the effect of preform preheating time on the interfacial region. The layer thicknesses of each reaction phase were measured in at least 10 different zones for each specimen and the average results were given in the Fig. 12. θ-Fe4Al13 phase thickness was increased with increasing preheating time while η-Fe2Al5 phase thickness was begun to decrease after showing a slight increase in 10 min

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preheated sample. 60 min preheated specimen did not include η-Fe2Al5 layer because of abovementioned reasons related to CrO2 formation. Another point to note was that the standard deviation of θ-Fe4Al13 thickness was considerably increased with increasing preheating time. It was the proof that the layer continuity of θ-Fe4Al13 phase was broken and the amount of discontinuous

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growth was increased with increasing preheating time.

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Figure 12. Preheating time dependent layer thickness values of θ-Fe4Al13 and η-Fe2Al5 phases After microstructural characterization, wear tests were conducted on the fabricated composites. For only this section of the complete work, nonreinforced A356 cast alloy was also examined to demonstrate reinforcement effect on wear behavior. Fig. 13 shows worn surfaces of A356 alloy in different magnifications. The worn surface of unreinforced alloy consists of extensive grooves parallel to the sliding direction, plastic flow marks, ridges and craters with miscellaneous dimensions. The existence of long continuous grooves approves the predominance of abrasive wear caused by harder asperities of the Al2O3 counterpart that plough the material surface. Likewise, ridges and craters indicate that adhesive wear and plastic deformation mechanisms are also in charge. Detached materials from the softer Al matrix were transferred to harder Al2O3 ball during test and either remained on ball or attached to matrix again. Another observation on Fig. 13a is that the dark

ACCEPTED MANUSCRIPT surfaces were locally covered by white thin layers which are indicative characteristics of oxidative wear. Frictional heating during sliding was the main reason of the oxidation. For Al-Si alloy like A356, oxidation did not play a destructive role under dry sliding conditions because oxide film prevents ball-disc contact and reduces adhesive wear [28]. Delamination wear was also occurred in Al matrix as shown in Fig. 13b. Sub-surface cracks, which may be whether pre-existed or newly formed due to

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and plastic deformation rates of A356 alloy were increased.

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generated shear strain during wear test, propagated and joined to wear process. Accordingly, wear

Figure 13. SEM images of worn surfaces of A356 alloy in (a) lower and (b) higher magnifications SEM images of worn surfaces of produced composites were given in Fig. 14. The first thing that

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seems pretty clear is that the widths of wear tracks for all reinforced samples were narrower than that of unreinforced A356 matrix. It can be obviously seen that 304 SS reinforcing provides decrease in wear rate. In general, typical smooth grooves (abrasion), flow marks (plastic deformation) and

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ridges (adhesion) were observed in all composite samples. θ-Fe4Al13 layer of nonpreheated specimen (Fig. 14b) was exposed to localized crack initiation and propagation during sliding. Particles from this cracked layer participated in wear process and these harder asperities damaged composite structure together with Al2O3 ball. Limited amount of delamination was obtained at this nonpreheated sample, especially in interfacial region. Shear strain generated by sliding caused crack propagation and propagated cracks caused delamination between matrix and reinforcement.

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Figure 14. SEM images of worn surfaces of 304 SS reinforced A356 matrix composites with (a-b) 0 min, (c-d) 10 min, (e-f) 30 min and (g-h) 60 min preform preheating

ACCEPTED MANUSCRIPT Severe wear of nonpreheated composite was changed to mild wear by preheating preforms for 10 and 30 min before casting. Fig. 14d shows that movement of sliding tracks in Al matrix was stopped by a reinforcement phase. Debris carried via wear tracks along the composite were accumulated in interfacial region, mostly in reinforcement side. Wider grooves and flow marks were attained in Al side of 30 min preheated sample than 10 min preheated one. However, double-layered interfacial

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phase of this specimen acted as a barrier to adhesive tracks. Only abrasive wear was occurred in reinforcement and interlayer sides while matrix side was deformed under both abrasive and adhesive conditions. It can be concluded that it is not enough to obtain interfacial layers including θFe4Al13 and η-Fe2Al5 phases to achieve good wear resistance, it is an obligation to produce these

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layers with strong bonding.

More debris can be seen in wear tracks of 60 min preheated samples (Fig. 14h) while the others have relatively uniform tracks with admissible amounts of wear debris. In this sample, long preheating

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time caused more oxidation at interface and CrO2 phase was observed instead of η-Fe2Al5. Because of the repeated ploughing of counterpart, oxide film was fatigued and delaminated from Al matrix, together with θ phase. Similar to nonpreheated sample, released oxide asperities joined to sliding

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tracks and increased wear damage caused by adhesion.

Figure 15. Sliding distance dependent friction coefficient values of produced samples The variation of coefficient of friction recorded by tribometer was given in Fig. 15. Friction coefficient of reference A356 alloy was higher than its composites reinforced with 304 SS. After approximately 80 m sliding distance, oxidation caused by frictional heating provided slight decrease in friction

ACCEPTED MANUSCRIPT coefficient by acting as a barrier layer. The minimum friction coefficient values were obtained in 10 and 30 min preheated samples. 10 min preheated sample was more resistant to wear during the first half of the test while 30 min preheated one showed better performance against wear for longer sliding distances. The friction coefficient of 60 min preheated sample was started to rise during the second half of the test and its value at 100 m distance was even higher than reference A356 sample.

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CrO2 film cracking during sliding may be the reason of this instantaneous increasing.

Figure 16. The variation of wear rate of produced samples Wear rates of manufactured composites were shown as a logarithmic scale in Fig. 16. Wear rates

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were calculated using WR=V/(L.P) formula where WR is wear rate, V is volume loss, L is sliding distance and P is applied load. Volume losses were derived from both analyzing wear tracks by using

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microscopy techniques and dividing calculated mass loss by density. The best results were obtained in 10 and 30 min preheated samples while the worst results were attained in reference Al alloy by a long way. These results were clarified by SEM images of worn surfaces. Nonpreheated sample was more resistant to wear than 60 min preheated sample. It was inferred that excessive preform preheating adversely affected on dry sliding wear resistance while sufficient one led to enhance tribological features. 4. CONCLUSION In this study, 304 SS reinforced A356 alloy matrix composites with varied preform preheating times were successfully manufactured by melt infiltration casting in an economic way. Double-layered intermetallic compounds including θ-Fe4Al13 and η-Fe2Al5 phases were obtained at interface owing to

ACCEPTED MANUSCRIPT long solid/liquid interaction, except for 60 min preheated sample. CrO2 layer was formed because of the excessive heat exposure in this specimen and this oxide film restrained η-Fe2Al5 formation. Worn surfaces of nonreinforced A356 alloy exhibit severe wear by comparison with its composite structures. Its worn surfaces were under the thumb of abrasive, adhesive, oxidative and delamination wear mechanisms. Reinforcing improved wear resistance of all composites while

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preheating affected wear behavior in different ways. It was understood that there is an optimum value for the preheating time. Preheated samples showed better improvement than nonpreheated sample until time was up to 60 min. As a consequence, it was concluded that sufficient preheating improved wear resistance of 304 SS reinforced A356 matrix composite under dry sliding conditions.

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Preheating preforms provide avoiding rapid solidification during melt infiltration casting Preheating with sufficient time improved wear resistance of composite due to well bonding Zero or excessive preheating times caused crack initiation and propagation at interface There is an optimum value for preheating time which varies from 10 to 30 min.

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