Casting temperature dependent wear and corrosion behavior of 304 stainless steel reinforced A356 aluminium matrix bimetal composites fabricated by vacuum-assisted melt infiltration casting

Journal Pre-proof Casting temperature dependent wear and corrosion behavior of 304 stainless steel reinforced A356 aluminium matrix bimetal composites fabricated by vacuum-assisted melt infiltration casting Ridvan Gecu, Ahmet Karaaslan PII:

S0043-1648(19)31196-2

DOI:

https://doi.org/10.1016/j.wear.2020.203183

Reference:

WEA 203183

To appear in:

Wear

Received Date: 1 August 2019 Revised Date:

23 December 2019

Accepted Date: 1 January 2020

Please cite this article as: R. Gecu, A. Karaaslan, Casting temperature dependent wear and corrosion behavior of 304 stainless steel reinforced A356 aluminium matrix bimetal composites fabricated by vacuum-assisted melt infiltration casting, Wear (2020), doi: https://doi.org/10.1016/j.wear.2020.203183. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2020 Published by Elsevier B.V.

Casting Temperature Dependent Wear and Corrosion Behavior of 304 Stainless Steel Reinforced A356 Aluminium Matrix Bimetal Composites Fabricated by VacuumAssisted Melt Infiltration Casting Ridvan Gecu1*, Ahmet Karaaslan1 *[email protected] 1

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

ABSTRACT This study aims to determine optimum casting temperature to achieve maximum wear and corrosion resistance in 304 stainless steel (SS) reinforced A356 matrix bimetal composites, which have been developed for use where conventional ceramic reinforced metal matrix composites have been used. Vacuum-assisted melt infiltration casting technique was used to manufacture bimetal composites at the various casting temperatures from 630 ˚C to 880 ˚C. The composite samples were subjected to dry sliding ball-on-disc tests using Al2O3 ball and immersion test in 3.5% NaCl solution for 7 to 21 days. θ (Fe4Al13) and η (Fe2Al5) phases developed at A356/304 interfaces and thickened with increasing temperature. Abrasion and adhesion as predominant wear mechanisms were observed in worn surface examinations, while pitting and galvanic corrosion occurred in corroded surfaces. The most suitable casting temperature was found to be 730 ˚C, considering the wear and corrosion properties of produced bimetal composites under certain conditions.

Keywords: Bimetal composite; melt infiltration casting; wear; corrosion; casting temperature; vacuum assistance

1. INTRODUCTION Ceramic reinforcements have been widely used in metal matrix composites (MMCs) which have been developed for applications where high strength, high wear resistance, good thermal stability, good corrosion resistance and low density are required. Aluminium and its alloys stand out as matrices in MMCs due to their unique properties such as ease of production, low density, high specific strength and good corrosion resistance [1]. These superior properties have enabled Al matrix MMCs (AMCs) to be used in variety of industries like aerospace,

automotive, electronics and defense [2]. Although SiC, Al2O3, TiC, B4C, TiB2, graphite or CNTs have been commonly preferred in AMCs because of their relatively high modulus of elasticity and fracture toughness [3], ceramic reinforcements show poor wettability against Al alloys [4]. Undesired intermetallic phase formation at the interface is the another problem besides wettability, when it comes to the metal/ceramic contacts [5]. In addition, ceramic reinforcements cause decrease in ductility and toughness of the matrix material [6]. The possibility of eliminating such existing problems caused by ceramic reinforcements in the production of conventional MMCs has been ignored. These production problems, which have not been fully solved yet, require a different approach to the issue. Unlike ceramic/metal composites, metal/metal composites have enormous potential to be manufactured more ecofriendly and economically without causing loss of ductility and toughness [7]. Recently, steel-based reinforcements show themselves up to enhance mechanical features of AMCs. For example, steel reinforced AMC sheets have been used as laminated tubes in air cooling systems of thermal power plants owing to excellent combination of high strength and good thermal conductivity of the steel with high corrosion resistance of Al alloy [8]. Khorrami et al. [9] researched on steel reinforced AMCs and determined that well-bonding interface and superior tensile strength of the bimetal composites were achieved by intermetallic compound (IMC) consisting of θ (Fe4Al13) and η (Fe2Al5) phases formed at the interface. Baron et al. [10] studied the IMC layer in AMCs and reported that excessive amounts of the IMC layer caused decrease in mechanical strength and there should be an upper limit for IMC volume fraction not to weaken the composite strength. Bhagat [11] worked on stainless steel (SS) reinforced AMCs and claimed that the optimum thickness of the IMC layer between SS and Al should be about 3 µm. In the following years, contrarily, Ozaki and Kutsuna [12] proved that if the IMC layer thickness was less than 10 µm, composite specimen was failed in tensile test because of the weak bond strength. In the authors’ previous work [13], the total thickness of the IMC layer consisting of θ and η phases reached up to 45 µm and increasing well-bonded IMC thickness provided increase in mechanical and wear properties. However, in following study [14], the total IMC thickness was in excess of 45 µm and consequently, tribo-mechanical features began to decrease in the composite specimen. When two different alloys are combined, it can be concluded that the interfacial reactions are very important to understand the composite features and they must be studied attentively.

On the other side, Selvakumar et al. [15] manufactured SS reinforced AMC without IMC formation at the interface. This research demonstrated the increase in composite ductility under favor of metal/metal interaction, by comparison with conventional TiC reinforced AMC in Thangarasu et al.’s [16] study. It was a particularly significant conclusion to show the potential of metal reinforcements to increase toughness and ductility, as well as ultimate tensile strength (UTS). Besides interface characterization, detailed microstructural, mechanical, wear and corrosion investigations are required to discover bimetal composites fully. Although there have been numerous studies on wear and corrosion behavior of the conventional ceramic reinforced AMCs, when the reinforcement type is metal, the existing literature is extremely insufficient. Only few working groups [17,18] have studied on AMCs reinforced with 10% steel at maximum. These researchers reported that the hardness, UTS and wear resistance values of the bimetal composites increased with increasing steel volume fraction. No studies have been made or suggested for higher volume ratio of the steel. Moreover, these studies are like technical reports and do not contain any information about the wear mechanisms of the examined composites. Further wear investigations are required to find out relationship between the interfacial reactions and the wear mechanisms. Generally, pitting, galvanic corrosion, stress corrosion cracking and fatigue corrosion types were observed in the conventional ceramic reinforced AMCs [19]. In bimetal composites, the predominant mechanism is galvanic corrosion due to wide differences between standard electrode potentials of the composite components. When the two metals are electrically bonded in conductive environment, the potential difference results in electron flux, and the metal with a more negative potential is preferably corroded [20]. Mandal et al. [21] stated that due to the formation of θ phase at the interface, galvanic cell was established between θ and Al matrix, instead of the matrix and the reinforcement phases. Depending upon the galvanic cell, it can be inferred that IMCs may have positive or negative effects on the corrosion resistance. The relationship between the interfacial reactions and the corrosion behavior in the bimetal composites should also be explained. Bimetal composites can be most properly produced by various manufacturing techniques that consist of molten metal routes, given manufacturing costs. Vacuum-assisted melt infiltration casting (VAMIC), as a molten metal process, allows high volume fraction of reinforcement phase, which participates in the composite production as a porous monoblock preform [22]. This technique uses plaster mould as a casting environment in order to maintain the

temperature before solidification as much as possible. When molten Al alloy maintains its liquid state for longer time, narrow preform cavities of the reinforcement phase can be easily infiltrated under vacuum atmosphere. Considering limited solidification time and high volume fraction together, the selection of the matrix and the reinforcement becomes crucial. When the mutual effects of wear and corrosion features were evaluated, A356 Al alloy was selected as the matrix because of its remarkable fluidity properties, while 304 SS was chosen as the reinforcement due to its excellent corrosion resistance [23–25]. In this study, A356 Al matrix bimetal composites reinforced with 304 SS were fabricated by VAMIC process at different casting temperatures varied from 630 to 880 ˚C with the interval of 50 ˚C. The influences of casting temperatures on interface, wear and corrosion features of the composites under certain conditions were determined.

2. EXPERIMENTAL PROCEDURE 2.1. Materials 304 SS reinforcement participated in the bimetal composite as sawdust form, in order to make environmentally friendly and economic production come true. 304 SS sawdusts, obtained from turning machine, were assembled by mechanical pressure to create porous monoblock preform, which was infiltrated by A356 melt via further VAMIC process. Chemical compositions of A356 and 304 SS alloys as the composite components are given in Table 1. Table 1. Chemical compositions of A356 Al and 304 SS 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.

-

2.2. VAMIC Process Bimetal composite production was carried out by VAMIC process which consists of preform preparing, mould making and infiltration stages, respectively. 2. 2. 1. Preform Preparing The porous SS preforms with dimensions of 20 mm in diameter and 10 mm in height were attained by gathering sawdusts together under 220 MPa pressure. This value was determined

by trial and error method. When the mechanical pressure was 220 MPa, obtained SS preform comprises 50% porosity. 2. 2. 2. Mould Making Commercial gypsum-bonded investment powder (Eurovest), which comprises silica, gypsum and various chemical additives, was used in the mould making step. While silica in the form of quartz and cristobalite regulates thermal expansion of the mould, the other additives are effective on hardening time, viscosity and strength. Castable plaster slurry was prepared by mixing the investment powder with water. Excessive use of the water in the mixture reduces the strength of the mould, while the use of excessive powder reduces its surface quality. In addition, long mixing times cause agglomeration and difficult mechanical agitation [26]. Due to the above-mentioned reasons, the water/powder ratio and the mixing time were determined as 40% and 3 minutes, respectively. A wax model with dimensions of 25 mm in diameter and 50 mm in height was prepared by pouring molten wax into a cylindrical plastic mould. The model was fixed to the mould base and settled into perforated SS flask. The reason for the perforation is to give the mould ability to take vacuum for infiltration. The prepared slurry was poured to the flask under vibration and the model was completely covered. After hardening of the slurry in an undisturbed condition, dewaxing process was carried out at 110 °C for 1 hour to dissipate wax pattern and generate mould cavity in plaster mould. The dewaxed mould was gradually heated up to 700 °C to dehydrate the mould and eliminate residual C from the removed wax. The mould was firstly heated to 250 °C and held at this temperature for 1 hour to complete αβ transformation of cristobalite. After that, the mould was heated to 550 °C and held at this temperature for 1 hour. While the mould was kept at 550 °C, it was aimed to carry out α-β transformation of the quartz, which causes significant volumetric expansion in the mould. When the temperature reached 700 °C, residual C combined with O2 in the air to form CO and CO2 gases and left the system completely. 2. 2. 3. Infiltration The VAMIC process is illustrated in Fig. 1. 304 SS preform was settled in the mould 10 min before the heat treatment of the mould was finished. The aim of the preform preheating for 10 min was to increase solid/liquid interaction time and achieve well-bonding at the interface. When the heating regime of the mould was completed, the perforated flask was immediately placed in a vacuum chamber. Gasket was interlaid between the flask and the chamber to

provide impermeability during casting. The A356 melt was cast into pouring basin and infiltrated the 304 SS preform cavities under -105 Pa vacuum pressure. When the melt was completely solidified, 50% 304 SS reinforced A356 matrix bimetal composite production was done. With the intent of studying the effects on bimetal composite properties, the casting temperatures of the molten A356 alloy was changed from 630 to 880 °C with the interval of 50 °C.

Figure 1. Illustration of VAMIC technique; (a) general view, b) cross-section view 2.3. Composite Characterization 2. 3. 1. Phase Determination and Microstructure Formed phases in the composites structure were determined by X-ray diffraction (XRD, Philips PW 3710) using CuKα radiation at 40 kV voltage and 40 mA current over 2θ range of 10-90°. Obtained XRD patterns were classified according to the The Joint Committee on Powder Diffraction Standards (JCPDS) database. The identified phases were given with the JCPDS numbers. Microstructural analysis of ground and mirror-polished cross-sections of the composite samples was performed by scanning electron microscope (SEM, Hitachi SU3500 T2). The thickness of the formed IMC layers at the interface were measured by image analysis software via optical microscope (OM, Nikon Eclipse MA100). Composition and distribution of elements in the formed phases were analyzed by energy-dispersive X-ray spectroscopy (EDS, Oxford XACT). 2. 3. 2. Hardness Measurements

The overall hardness of the bimetal composites was measured by Brinell hardness tester with load of 62.5 kg and tip diameter of 2.5 mm in accordance with BS EN ISO 6506-3:2014 standard. The results were given as the mean value of 5 measurements taken from different regions for each sample. Besides the overall hardness, the individual hardness of the formed phases in the composite structure was also calculated by nanoindentation test device (CSM NHT, SN06-177). Random measurements were taken from 10 different regions for each phase and Vickers hardness results were given as the mean value of these measurements. The nanoindentation test device uses the Oliver and Pharr nonlinear curve fitting method to measure hardness [27]. 2. 3. 3. Wear Tests The specimens were subjected to unlubricated ball-on-disc wear tests by using 6-mmdiameter Al2O3 ball as a counterface. The specifications of the counter material are given in Table 2. The main reason for the selection of Al2O3 ball is that its hardness is higher than any of formed phases in composite structure. By using Al2O3 ball under dry sliding conditions, it is aimed to explore relatively severe wear conditions. Table 2. Specifications of Al2O3 ball as a counter material Characteristic

Magnitude

Characteristic

Magnitude

Purity (%)

99.5

Hardness (HV10)

1410

Size (mm)

Φ6

Young’s modulus (GPa)

370

Density (g/cm3)

3.86

Compressive strength (MPa)

2600

Water absorption (%)

< 0.01

White

Color

Other test parameters were chosen to represent the above-mentioned severe wear conditions. 10 N load was applied to polished composite surfaces during 100 m sliding distance at 0.1 m/s sliding velocity. The arithmetic roughness (Ra) of the polished surfaces was measured as 0.45 µm. During sliding, ambient temperature and relative humidity of the air were kept at 24 °C and 40%, respectively. Three wear tests were performed for each sample and the average specific wear rate results were given with the standard deviations. =2



⁄2

−

⁄4 4

−

(1)

equation was used to calculate volume losses of the bimetal composites, where V is the volume loss, R is the wear track radius, r is the ball radius and d is the wear track width, in accordance with ASTM G99-17. By use of the obtained volume losses, the specific wear rates were measured by following equation, where WR is the wear rate, L is the sliding distance and P is the applied load. WR =

V LP

(2)

Within the scope of the wear characterization, coefficient of friction (CoF) values were also recorded by the tribometer. The CoF values depending on the sliding distance were given as the average of the three test results for the each composite specimen. When the dry sliding tests were completed, the worn surfaces of the all bimetal composite samples were examined by SEM to determine wear mechanisms. 2. 3. 4. Immersion Tests The samples were immersed in aerated 3.5% NaCl solution and held between 7 and 21 days, to determine corrosion mechanisms that occurred. At the end of the tests, corrosion products were mechanically removed from the composite surface. XRD analysis was carried out to identify the corrosion products, whereas the micrographs of the corroded surfaces were taken by OM. Corrosion rates were calculated by using

(8.76 x10 4 xW ) CR = ( AxTxD) (3) equation, where CR is the corrosion rate, W is the mass loss, A is the surface area, T is the immersion time and D is the density of the specimen.

3. RESULTS AND DISCUSSION 3.1. XRD Analysis Fig. 2 shows X-ray diffractograms of the bimetal composites cast at different temperatures. The JCPDS numbers of the identified phases are reported in Table 3. Al, Fe, θ and η phases formed in the composite structure for the both temperatures at 680 and 880 °C. As the temperature increased, the intensities of the θ and the η diffractograms increased due to longer interaction time between the liquid matrix and the solid reinforcement phases.

According to Fe-Al phase diagram [28], when Fe atoms diffuse into molten Al for a sufficient time, the first reaction phase is θ, which is the one of the products of the eutectic reaction at 654 °C. When the interaction time is long enough, η, ζ (FeAl2) and α’ (FeAl) phases can also form. Considering the phase diagram and the XRD results together, it can be inferred that the solid/liquid interaction time is long enough to form the θ and the η phases, while it is short enough not to allow the formation of the ζ and α’phases.

Figure 2. X-ray diffractograms of the 304 SS reinforced A356 matrix bimetal composite produced at 680 and 880 °C casting temperature

Table 3. JCPDS numbers of the identified phases in Fig. 2 Phase

JCPDS number

Al

00-004-0787

Fe

00-001-1267

θ (Fe4Al13)

00-029-0042

η (Fe2Al5)

00-047-1435

3.2. Microstructure

Fig. 3 shows SEM images of the 304 SS reinforced A356 Al matrix bimetal composites produced by the VAMIC technique. Dark regions point out the matrix phase while light regions imply the 304 SS reinforcement. Numerous discontinuities at the 304/A356 interfaces are marked in Fig. 3a and Fig. 3b. These defects occurred because of the rapid solidification of the A356 alloy at temperatures close to its melting temperature.

The rapid solidification was prevented by increased temperature. Double-layered continuous IMC layer formed at the all specimens manufactured above 730 °C. Especially the sample cast at 730 °C has an almost perfect microstructure that does not involve any discontinuity at the interface or shrinkage cavity in the matrix. As the temperature increased, huge difference between the temperatures of the liquid and the solid phases caused increase in thickness of the θ and the η layers. This thermal gradient also caused the shrinkage cavity formation in the matrix phase because of increased thermal stress between the metal pairs during casting. According to the microstructural analysis, both low and high temperatures caused the casting defects. The temperature should be chosen as low as to prevent the shrinkage cavities and as high as not to allow the interface discontinuities. Larger scale SEM images of the bimetal composites fabricated at varied casting temperatures are given in Fig. 4. The first thing that stands out in the images is that the thickness of the θ and the η phases increased as the temperature increased. The disintegration caused by mechanically weak bonding between the matrix and the reinforcement is clearly seen at lower temperatures (Fig. 4a and Fig. 4b). Very rapid cooling restrained to achieve well-bonding at these temperatures. As the temperature increased, strong mechanical and metallurgical bonding was obtained at the interface (Fig. 4c and Fig. 4d). However, further increase in the temperature caused discontinuous growth of the θ and the η layers. The brittleness of these reaction phases increased as their thickness increased. Consequently, the specimens cast at 830 and 880 °C were damaged without exposed to any external forces (Fig. 4e and Fig. 4f). Considering both Fig. 3 and Fig. 4, it was concluded that the optimum casting temperature should be in the range of 730-780 °C.

Figure 3. SEM images of the 304 SS reinforced A356 matrix bimetal composite cast at (a) 630, (b) 680, (c) 730, (d) 780, (e) 830 and (f) 880 °C (white arrows indicate the interface discontinuities)

Figure 4. Larger scale SEM images of the 304 SS reinforced A356 matrix bimetal composite cast at (a) 630, (b) 680, (c) 730, (d) 780, (e) 830 and (f) 880 °C 3.3. EDS Analysis Fig. 5a shows SEM image and EDS spots of the 304 SS reinforced A356 matrix bimetal composite cast at 730 °C. The formed phases that previously determined by the XRD analysis are clearly seen in different colors. From up to down, these phases are the matrix, the doublelayered IMC consisting of the θ and the η phases, and the reinforcement, respectively. The results of the EDS analysis are reported in Table 4. Spot 1 indicates the A356 matrix where αAl and eutectic (α-Al + Si) phases exist together, while spot 8 implies 304 SS reinforcement undoubtedly. Both the θ and the η phases are analyzed at three different spots. According to

the EDS results of the IMC layer, the alloying elements are homogenously distributed along both the θ (spots 2-4) and the η (spots 5-7) phases. The remarkable point is that Si atoms tend to accumulate into the θ layer rather than the η. It can be inferred that the formation of the η phase acts as a barrier to diffusion of the Si atoms.

Figure 5. (a) SEM image and EDS spots of the 304 SS reinforced A356 matrix bimetal composite cast at 730 °C and (b) elemental content variation as a function of distance

Table 4. EDS analysis results (wt.%) in Fig. 5 Spot

Al

Si

Fe

Cr

Mo

Ni

C

1

84.21

10.96

-

-

-

-

4.83

2

62.87

8.60

18.02

6.78

1.31

-

2.42

3

61.36

8.46

19.49

7.32

1.31

-

2.06

4

59.94

9.03

21.28

6.33

1.92

0.38

1.12

5

55.70

-

28.45

9.08

2.41

1.82

2.54

6

56.64

-

29.61

8.23

1.88

2.01

1.63

7

56.17

-

29.10

10.05

2.23

1.56

0.89

8

-

-

69.52

17.87

2.47

5.64

4.50

Elemental content variation derived from the EDS results given in Table 4 is illustrated in Fig. 5b to clarify the distribution of the alloying elements along the line marked in Fig. 5a. From

the matrix side to the reinforcement side, the amounts of Al and Si decrease while the quantities of Fe and Cr increase. 3.4. Thickness and Hardness of the IMC Layer The thickness values of the IMC layer depending upon the casting temperature of the 304 SS reinforced A356 matrix bimetal composites are demonstrated in Fig. 6. The thicknesses were measured at 10 random regions for each reaction phase and each specimen by image analyzer software via OM.

Figure 6. Casting temperature dependent thickness of the IMC layer formed at the interface of the 304 SS reinforced A356 matrix bimetal composite As the casting temperature increases, the thickness of the θ layer shows a continuous increase, while the thickness of the η increases at first and then decreases. The reason for the continuous increase in the θ thickness is increased diffusivity with the increasing temperature. The increase in the diffusivity means an increase in kinetic energy, which lead to more spontaneous spreading of the alloying elements. Due to increase in the kinetic energy, the diffusion occurs more quickly and consequently the θ thickness increases. Considering huge standard deviation at 830 and 880 °C, it can be said that, in cases where the temperature is more than 780 °C, the θ layer grows discontinuously. At some regions, the θ thickness reached almost 120 µm, although the average thickness varied from 40 to 65 µm at these temperatures.

According to the Fe-Al phase diagram, besides the eutectic reaction, θ can be formed as a result of the peritectic reaction between liquid Al and η. It was thought that the decrease in η thickness above 780 °C was due to the fact that a part of the formed η comes together with the Al melt to form the θ phase. When the solid/liquid interaction is kept longer, the growth rate of the η slows down and the transformation tendency from η to θ stands out. Because of the above-mentioned relationship between the liquid Al and the IMC layer, the thickness of the η layer began to decrease when the casting temperature exceeded 780 °C. As previously explained, the IMC layer should be at least 10 µm in order to avoid mechanical damage at the interface [12]. All of the manufactured composites contained the IMC layer with the thickness above the critical threshold. It was also known that when the total thickness of the θ and the η phases exceeded 45 µm, the composite properties weaken [14]. In the light of such information, it was determined that the casting temperature should not exceed 730 °C to achieve the optimum thickness. The effect of the casting temperature on the hardness of the IMC layer is given in Fig. 7. The hardness of both the θ and the η increased with the increase in the temperature up to 780 °C. The θ and the η lattices were distorted by the diffusion of the alloying elements during the production. The increase in the casting temperature provides inhibiting rapid solidification and the alloying elements find an opportunity to diffuse in the lattices for longer time. As a result of disrupted coherency in the lattices due to the increased diffusivity of the alloying elements, the hardness values of the θ and the η layers increased. However, the further increase above 780 °C caused a decline in the hardness owing to the longer interaction time before solidification. Due to the delay of the solidification, the transformation from the η to the θ resulted in the elimination of mismatches caused by the diffusion of the alloying elements into the η lattice. It was thought that the hardness of the η phase decreased because of the removal of these lattice mismatches. On the other hand, the decrease in the θ hardness with the increasing temperature can be explained by the increased brittleness, which may cause an underestimation of the hardness due to crack initiation and propagation at the interface. During loading, the nucleated cracks may rapidly propagate, which lead to a sharp increase in the penetration depth of the indenter resulting in a miscalculation of the hardness.

Figure 7. Casting temperature dependent hardness of the IMC layer formed at the interface of the 304 SS reinforced A356 matrix bimetal composite

Figure 8. Thickness-hardness relationship of the IMC layer depending upon the casting temperature Considering the all temperatures, the average hardness of the θ varied from 807 to 903 HV0.01 while the average η hardness varied between 1027 and 1208 HV0.01. Although the θ hardness is consistent with the study of Matysik et al. [29], the hardness of the η layer is far out of the given range. Owing to the material pile-up around the contact impression, the hardness results can be over-estimated in the Oliver and Pharr nonlinear curve fitting method

[30]. To reveal the correlation between the thickness and the hardness of the IMC layer, thicknesshardness relationship depending upon the casting temperature is demonstrated in Fig. 8. Except the two samples cast at 830 and 880 °C, the rest of the bimetal composites showed the linear relationship between the thickness and the hardness for both the θ and the η layers. As mentioned before, the specimens produced at these temperatures became brittle with the increasing thickness, resulting in the erroneous measurements. 3.5. Hardness of the Matrix and the Reinforcement Table 5 gives the Vickers hardness values of the matrix and the reinforcement phases, depending upon the casting temperatures. The hardness values varied from 49 to 71 HV0.01 for the A356 matrix while varied between 233 to 245 HV0.01 for the 304 SS reinforcement. It was understood from the results that the hardness of the 304 SS did not significantly depending on the temperature, while the hardness of the matrix decreased with the increasing temperature due to the grain growth in the α-Al phase. Especially in the sample cast at 880 °C, the hardness declined to 49 HV0.01 because of the excessive grain growth.

Table 5. Casting temperature dependent hardness of the matrix and the reinforcement phases Casting temperature (°C)

Hardness (HV0.01) Matrix

Reinforcement

630

71

245

680

69

241

730

67

239

780

61

241

830

52

233

880

49

238

3.6. Worn Surface Examinations Fig. 9 shows worn surface SEM images of the 304 SS reinforced A356 matrix bimetal composites. The worn surfaces of the bimetal composites consist of fragmented particles, wide grooves parallel to the sliding direction, plastic flow marks, delaminated interfaces, ridges and craters of all sizes. The existence of the continuous grooves and the detached

regions proves that the abrasive wear is a predominant mechanism. During abrasive wear, the material loss can be occurred by microcutting or microploughing mechanisms. The microcutting mechanism results in more material loss than the microploughing, which is based on the displacement of the material by generating boundaries. The micro-cutting causes particle fragmentation while the microploughing generally results in the formation of the grooves [31]. According to the worn surface SEM images, the abrasive wear occurred as the microploughing up to 730 °C. The microploughing transformed into the microcutting above this temperature. The detached particles in black color in Fig. 9e and Fig. 9f are resulted from the microcutting.

Figure 9. Worn surface SEM images of the 304 SS reinforced A356 matrix bimetal

composites produced at (a) 630, (b) 680, (c) 730, (d) 780, (e) 830 and (f) 880 °C

In previous studies [32,33], it was determined that the high wear-resistant material was deformed under the microploughing, whereas the low wear-resistant material was mostly disintegrated under the microcutting. In this case, it can be expected to obtain relatively low wear resistance especially in the bimetal composite samples fabricated at 830 and 880 °C, because the amount of the fragmented particles was higher than the other specimens. Besides the abrasive wear, the adhesive wear also damaged the bimetal composite samples. The presence of the ridges and the craters in Fig. 9 can be counted as an evidence of the adhesive wear. Fragmented particles, plastered to Al2O3 ball during sliding, either participated to the composite structure again or remained on the ball. This cycle made the ridges and the craters on the worn surfaces, respectively. Another remarkable point is that the amount of the delaminated regions at the interface, caused by an insufficient adhesion, decreased with the increasing temperature up to 780 °C. However, especially at 830 and 880 °C, the amount of the interface delamination dramatically increased due to the increased brittleness of the IMC layer. In areas exposed to delamination in the samples manufactured at higher temperatures, wear debris density increased. It was understood from the SEM images that the debris densities vary according to the casting temperature. The debris formation can be resulted from abrasion, adhesion, oxidation or delamination and its impact on the composite properties may be positive or negative depending upon the formation type [34]. Wear debris SEM images of the bimetal composites produced by VAMIC technique are given in Fig. 10. The size of the debris decreased as the casting temperature increased up to 730 °C and tended to increase again above this temperature. When the density was relatively low, small and near-spherical debris was formed, whereas large and spherical debris was generated due to agglomeration of the small debris, when the debris density was relatively high. The agglomeration occurs due to thermo-mechanical fusion of the debris formed under repeated loads during sliding [35]. Singh and Chauhan [36] reported that wear rate decreases as the wear debris size decreases. In the same study, it was stated that high wear-resistant composite production can be achieved by a coaxial small-sized debris formation with a stable tribolayer on the wear surface. Thus, it can be expected that the bimetal composites produced at 730 and 780 °C, having small, near-spherical debris, will have higher wear resistance than the other

samples.

When the hardness results (Fig. 7) and the wear debris images (Fig. 10) were evaluated together, it can be seen that the amount of the formed debris decreases as the hardness increases at the interface. These findings obtained from the wear debris SEM images are consistent with those of Lemm et al. [37].

Figure 10. SEM images of wear debris of the 304 SS reinforced A356 matrix bimetal composites produced at (a) 630, (b) 680, (c) 730, (d) 780, (e) 830 and (f) 880 °C Worn surface images of the Al2O3 balls against the bimetal composites manufactured at

varying temperatures are given in Fig. 11. It was clearly seen that the amount of the detached particles adhered to the ball is the highest for the specimen produced at 880 °C due to its increased brittleness-induced interface delamination. Although the other worn surface images look similar for the samples cast at 680 and 780 °C, the counterpart contains more attached particles from the bimetal composite fabricated at 680 °C because the adhesive wear of this sample is higher than that of the sample produced at 780 °C.

Figure 11. Worn surface images of the Al2O3 balls against the composite specimens produced at (a) 680, (b) 780 and (c) 880 °C 3.7. Specific Wear Rate and Coefficient of Friction Fig. 12 demonstrates the specific wear rate results of the bimetal composites as a function of

casting temperature. The increase in the casting temperature improved the wear resistance up to 730 °C. Above this temperature, however, it had a detrimental effect on the wear resistance. The increase in the specific wear rate is caused by the weak bonding at the interface in the samples cast at lower temperatures, whereas, in the specimens produced at higher temperatures, it is mainly related to the fragmentation of the discontinuous growing brittle phases at the interface. On the other hand, the double-layered continuous IMC layer provided maximum wear resistance in the specimen cast at 730 °C.

Figure 12. Specific wear rates of the 304 SS reinforced A356 matrix bimetal composites depending upon the casting temperature Within the scope of the tribological characterization, the CoF values were recorded by the tribometer during sliding process. Fig. 13 shows the sliding distance dependent CoF values of the 304 SS reinforced A356 matrix bimetal composites. It was obviously seen that the CoF values did not change with the sliding distance significantly. This steady-state of the CoF during sliding can be explained by the low “stick-slip effect” [38]. The oxidation of the contact surfaces at the beginning of the test hinders the “hold-release movement” and ensured that the samples maintain their steady-state throughout the entire process. The lowest CoF results were obtained in the specimens cast at 730 and 780 °C with average CoF of ~0.125 and ~0.133, respectively. The CoF values appear to be consistent with the specific wear rate results. The lowest CoF results were attained in the sample that has the maximum wear resistance compared to the others.

Figure 13. Sliding distance dependent CoF values of the 304 SS reinforced A356 matrix bimetal composites produced at various casting temperatures

Figure 14. Relationship between the specific wear rate and the layer hardness in the 304 SS reinforced A356 matrix bimetal composites depending upon the casting temperature The layer hardness results given in Fig. 7 and the specific wear rates shown in Fig. 12 were combined to reveal correlation between the hardness and the wear resistance. Fig. 14 shows the relationship between the specific wear rate and the layer hardness in the bimetal composites depending upon the casting temperature. It can be surely said that the increase in the hardness of the θ and the η phases provided decrease in the specific wear rates, provided that some exceptions (the results at 880 °C for θ and 830 °C for η) were ignored.

3.8. Corrosion Behavior Fig. 15 shows uncorroded and corroded regions of the bimetal composite sample after 7 days immersion in 3.5% NaCl solution. The left side represents liquid-tight coated region that was not affected by the solution. In the corroded part, a dissolution of the A356 matrix due to galvanic interaction occurred. No dissolution was observed in the other components of the composite, including 304, θ and η regions. It was also confirmed that the α-Al and the eutectic phases in the A356 matrix formed a galvanic cell, too. As a natural result of this formed cell, the dissolution in the eutectic phase was higher because of its larger heterogeneity than α-Al.

Figure 15. Uncorroded and corroded regions of the 304 SS reinforced A356 matrix bimetal composite after 7 days immersion in 3.5% NaCl solution The microstructures of the bimetal composites after 21 days immersion test in 3.5% NaCl solution are given in Fig. 16. The galvanic cell between the matrix and the reinforcement caused most severe damage to the specimen cast at 630 °C, whereas the least damage occurred in the sample produced at 730 °C. Especially in the composites produced at 830 and 880 °C, besides the A356 matrix, the SS surfaces have also become susceptible to the corrosion. At these temperatures, the diffusion of Cr atoms from SS side is at the maximum level, leading to the formation of Cr-poor regions in the reinforcement side when solidification was completed [39,40]. These Cr-poor regions have lost their SS properties and made the reinforcement vulnerable to the corrosion. The selective corrosion in the white zones shown in Fig. 16e and Fig. 16f has progressed through the entire reinforcement, with

deterioration of the passivity. The corrosion in SS part may also occurred because of the degradation of the protective Cr2O3 film in O2-containing atmospheres [41].

Figure 16. Microstructures after 21 days immersion of the 304 SS reinforced A356 matrix bimetal composites cast at (a) 630, (b) 680, (c) 730, (d) 780, (e) 830 and (f) 880 °C The sample produced at 880 °C was subjected to XRD analysis after 21 days immersion to determine the formed corrosion products. According to the XRD results given in Fig. 17, it was determined that the corrosion products in the composite structure are AlOOH (JCPDS; 00-033-0018) and FeOOH (JCPDS; 00-034-1266). While FeOOH is formed by either the oxidation or the precipitation mechanisms in aqueous solutions containing Fe+2 or Fe+3 ions

[42], for AlOOH formation, AlCl3 must first be formed [43]. Anodic and cathodic reactions occurring throughout the immersion are listed below. [44–46]. Fe  → Fe +2 + 2e −

(4)

O2 + 2 H 2 O + 4e −  → 4OH −

(5)

Fe +2 + 2OH −  → FeOOH + H + + e −

(6)

Al  → Al +3 + 3e −

(7)

Al +3 + 3Cl −  → AlCl3

(8)

AlCl3 + 2 H 2 O  → AlOOH + 3HCl

(9)

Figure 17. XRD analysis results of the corroded 304 SS reinforced A356 matrix bimetal composite cast at 880 °C Fe and Al were oxidized by the anodic reactions given in Eq. (4) and Eq. (7), while the cathodic reaction was carried out as described in Eq. (5). OH- ions, formed as a result of reduction of O2 at the cathode, reacted with Fe+2 ions in accordance with Eq. (6), and generated the corrosion product (FeOOH) seen in the form of white islets in Fig. 16e and Fig. 16f. Al+3 ions reacted with Cl- ions to form AlCl3 as a by-product as given in Eq. (8) and AlOOH was formed by the reaction of this by-product with H2O in the solution as described in Eq. (9). Cl- ions have just acted as catalysts to accelerate the corrosion and they were not included in the corrosion products [47].

Corrosion rates after immersion tests for 7, 14 and 21 days are given in Fig. 18. The best corrosion resistance was achieved in the sample produced at 730 °C for all immersion times. It was determined that the corrosion rate increased as the immersion time increased. Initially, the galvanic interaction caused rapid increase in the corrosion rate but then rate of the propagation of the corrosion in both the composite components slowed down due to their tendency to form passive films on their surfaces. The fact that the corrosion rates obtained at the end of 14 and 21 days is quite close to each other confirms this inference.

Figure 18. Corrosion rate results of the 304 SS reinforced A356 matrix bimetal composites as functions of the immersion time and the casting temperature

4. CONCLUSION 304 SS reinforced A356 matrix bimetal composites were successfully fabricated at varying temperatures from 630 to 880 °C with the interval of 50 °C by VAMIC, which is an environmentally friendly low cost preform-based process. The findings of the current study revealed the following conclusions:

•

Vacuum assistance allowed excellent infiltration of the molten A356 alloy during the operation for all casting temperatures, even at 630 °C. It was determined that the optimum casting temperature should be in the range of 730-780 °C to obstruct the shrinkage cavities in the matrix and the discontinuous growth of the IMC layer at the interface.

•

Double-layered IMC consisting of θ (Fe4Al13) and η (Fe2Al5) phases was obtained for all temperatures. While the θ thickness continuously increased as the casting temperature increased, the decrease in the η thickness when the temperature exceeded 780 °C was originated from the peritectic reaction between liquid Al and η to form θ.

•

It was determined that the predominant mechanism on the matrix side was adhesive wear whereas the predominant mechanism on the reinforcement side was abrasive wear under predetermined conditions. The abrasive wear occurred as the microploughing up to 730 °C, while the microcutting stepped in above this temperature.

•

θ and η phases enhanced the wear resistance of the bimetal composite when it was strongly bonded to the interface. However, when these interfacial phases were weakly bonded, they caused a dramatic increase in the wear damage by causing disintegration of themselves from the interface.

•

As a result of the immersion tests, no dissolution was observed in the 304, θ and η regions up to 780 °C. However, in the composites produced at higher temperatures, Cr-poor regions, formed in 304 SS side during casting, caused the deterioration of the passive film and make the reinforcement vulnerable to the selective corrosion.

•

Because of the galvanic interaction, A356 matrix acted as a protective cathode for the other components and dissolved. The eutectic region, where the heterogeneity is higher, dissolved more than the α-Al grains.

•

The best wear and corrosion resistance was achieved by the sample cast at 730 °C. IMC thickness, IMC hardness, wear rate, CoF and corrosion rate values of this specimen created the perfect combination of the wear- and corrosion-resistant composite structure under certain conditions.

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•

Double-layered IMC consisting of θ (Fe4Al13) and η (Fe2Al5) phases was obtained for all temperatures.

•

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•

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