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Article

Effect of Ti Addition on the Hot-Tearing Susceptibility of the AlSi5Cu2Mg Alloy

by
Marek Matejka
,
Dana Bolibruchová
and
Martina Sýkorová
*
Department of Technological Engineering, Faculty of Mechanical Engineering, University of Žilina, Univerzitná 8215/1, 010 26 Žilina, Slovakia
*
Author to whom correspondence should be addressed.
Metals 2024, 14(6), 703; https://doi.org/10.3390/met14060703
Submission received: 25 May 2024 / Revised: 10 June 2024 / Accepted: 12 June 2024 / Published: 14 June 2024
(This article belongs to the Special Issue Light Alloy and Its Application II)
Browse Figures
Versions Notes

Abstract

:
The aluminum alloy AlSi5Cu2Mg finds application in the production of high-stress cylinder head castings. The AlSi5Cu2Mg alloy is specific for its high susceptibility to hot tearing. One effective way to reduce the susceptibility of Al-Si-Cu-Mg alloys to hot tearing is by grain refining. The AlSi5Cu2Mg alloy is designed with a specific chemical composition that significantly limits the Ti content to a maximum of 0.03 wt.%. This limitation practically limits the use of standard Al-Ti-B-based refiners. The present work focuses on the investigation of the influence of graded Ti addition on the susceptibility of the AlSi5Cu2Mg alloy to hot tearing. The Ti addition was deliberately chosen beyond the manufacturer’s recommendation (0.1, 0.2, 0.3 wt.%). The solidification process of the experimental alloys with Ti addition was evaluated in this research. On the basis of the thermal analysis, it was shown that due to the addition of Ti, the solidification interval of the AlSi5Cu2Mg alloy increases. An increase in the solidification interval is often associated with an increase in the susceptibility to tearing. The susceptibility of the experimental alloys to hot tearing was evaluated qualitatively and quantitatively. Based on the quantitative and qualitative evaluation, it was shown that the addition of Ti reduces the susceptibility of the AlSi5Cu2Mg alloy to hot tearing. A positive refining effect of Ti on the primary α-(Al) phase was demonstrated by microstructural evaluation. Based on this research, it was shown that despite the increase in the solidification interval due to the addition of Ti, the susceptibility of the aluminum alloy to the formation of hot tears is reduced due to the better filling of the material in the interdendritic spaces.

Graphical Abstract

1. Introduction

Hot tearing is one of the major casting defects of the hypoeutectic aluminum alloy AlSi5Cu2Mg used for cylinder head castings [1]. Hot tearing, often referred to as hot cracking, occurs when a macrocrack forms during the solidification process [2]. Tear initiation occurs at temperatures near the solidus at locations of the lowest strength [3]. Such locations are primarily the grain boundaries, so the emerging flaw also runs along these boundaries, i.e., intercrystalline [4,5]. The cause of hot tearing is usually endogenous internal stress, most often due to the uneven cooling of the casting, inhibited shrinkage, or a change in density due to phase changes [6]. In addition, external factors that cause exogenous stresses, such as the resistance of the mold and core to shrinkage, also influence tear formation [7,8].
An equally important factor influencing tearing is the chemical composition of the material [9]. This affects the resulting susceptibility of the alloy to tearing, particularly in terms of the influence of individual chemical elements on the width of the solidification interval [10]. In the case of Al-Si-Cu-Mg alloys, copper has a significant influence on the susceptibility of the alloy to tearing, mainly through the influence of the width of the solidification interval [10,11]. In general, as the Cu content in the alloy increases, the solidification interval also increases, resulting in a decrease in the resistance of the alloy to tearing [12]. In general, the tearing susceptibility of Al alloys is directly proportional to the width of the solidification interval [10]. Studies have shown that one of the ways to reduce the susceptibility of the alloy to hot tearing is the grain refining of the Al-Si-Cu-Mg alloy [13].
Titanium is used for the grain refining of hypoeutectic Al-Si alloys. The grain refining of aluminum alloys is the process of deliberately introducing suitable elements or their alloys to refine the structure by increasing the number of crystallization seeds for the α-(Al) phase [14]. In the grain refining of aluminum alloys with titanium, Ti reacts with Al to form heterogeneous crystallization nuclei Al3Ti with a size of 10–20 µm [15]. As a result of the change in the solubility of Ti in aluminum, Al3TiII crystal nuclei are formed. Around the Al3Ti and Al3TiII crystal nuclei, a solid solution α-(Al) shell is formed, which acts as an active substrate for the growth of Al dendrites [14,15]. Al3Ti crystallization nuclei have a relatively high solubility in Al, which reduces the refining effect of Ti [15,16]. Studies have shown that a better refining effect of hypoeutectic Al-Si alloys is achieved by the combined effect of Ti and B [15,17]. Ti and B are used in the form of master alloys with different proportions of Ti (max. 5 wt.% Ti) and B (max. 1 wt.%). The principle of refining the Al-Si alloys with the Al-Ti-B master alloy is the formation of crystallization intermetallic phases based on AlB2, TiB2 or (Al,Ti)B2 with a size of 0.5–2 µm [18,19]. The crystallization nuclei are similar to Ti refining and consist of a core and a shell [20]. The core is formed by the (Al,Ti)B2 phase and the shell is formed by the TiAl3 phase, which acts as an active substrate for the growth of Al dendrites [18]. The grain refining of aluminum alloys leads to improved casting properties, mechanical properties and a reduction in susceptibility to hot tearing. Sigworth demonstrated that the addition of Ti leads to a reduction in the susceptibility of aluminum alloys to hot tearing [15] due to the better deposition of the melt in the interdendritic spaces [21]. Reducing the susceptibility of Al-Si-Cu-Mg alloys to hot tearing by means of grain refining can significantly expand the application range of high-susceptibility Al-Si-Cu-Mg alloys.
This paper focuses on the investigation of the effect of graded Ti addition on the hot-tearing susceptibility of the AlSi5Cu2Mg alloy. The non-standard AlSi5Cu2Mg alloy is used in the production of highly stressed cylinder head castings and is characterized by a high susceptibility to hot tearing. The AlSi5Cu2Mg aluminum alloy was designed by the manufacturer with a specific chemical composition compared to the conventional Al-Si-Cu-Mg alloys used for cylinder head castings. The AlSi5Cu2Mg alloy is characterized by a low Si content (5.0–6.5 wt.% Si). The AlSi5Cu2Mg alloy is similar to Al-Si alloys intended for forming regarding the Si content. The Si content of conventional Al-Si-Cu-Mg-based aluminum alloys intended for highly stressed cylinder head castings is in the range of 6.0–10.0 wt.% Si. The AlSi5Cu2Mg alloy is significantly limited by the Ti content (max. 0.03 wt.%). Optimum refinement of the structure of hypoeutectic aluminum alloys is achieved by adding 0.10 wt.% of Ti. The regulation of the chemical composition of the AlSi5Cu2Mg alloy set by the manufacturer, which significantly limits the Ti content in the melt (max. 0.03% by weight), significantly limits the use of standard Al-Ti and Al-Ti-B refiners. Despite the limitations imposed by the chemical composition regulation, the experimental work deliberately selected Ti additions (0.1, 0.2 and 0.3 wt.%) that exceeded the manufacturer’s recommendations. The main objective of the research was to determine the extent to which the excess amount of Ti affects the resulting susceptibility to hot tearing. This paper continues previous research that investigated the synergistic effect of Zr and Ti on the susceptibility of the AlSi5Cu2Mg alloy to hot tearing [22].

2. Materials and Methods

The hypoeutectic aluminum alloy AlSi5Cu2Mg (Nemak, Slovakia) was selected for this research. The AlSi5Cu2Mg alloy was designated as the reference alloy. The chemical composition of the reference alloy listed in Table 1 was determined using a calibrated SPECTROLAB S optical emission spectrometer (Spectro Analytical Instruments, Kleve, Germany).
The hypoeutectic aluminum alloy AlSi5Cu2Mg was melted in an electric resistance furnace LAC T15 (LAC, Židlochovice, Czech Republic) with an open atmosphere. The experimental alloys were obtained by refining the reference alloy AlSi5Cu2Mg with graded additions of Ti (0.1, 0.2 and 0.3 wt.% of Ti), while the Ti content was deliberately chosen to exceed the manufacturer’s recommendation (max. 0.03 wt.% Ti). Depending on the addition of Ti, the experimental alloys were designated Ti-1, Ti-2 and Ti-3. Ti was added to the melt in the form of the AlTi5B1 master alloy (Nemak, Slovakia) at a temperature of 770 °C ± 5 °C. The melt temperature was measured with an AHLBORN THERM 2420-1L thermometer (AHLBORN, Prague, Czech Republic) with a type K thermocouple (NiCr-Ni). The experimental alloys for the evaluation of the tear susceptibility were produced by gravity casting in an open atmosphere into a permanent mold. The mold was coated with graphite before the casting and the temperature of the mold was preheated to 150 °C ± 10 °C prior to the casting. The mold temperature was measured with a VOLTCRAFT IR 900-30S thermometer (VOLTCRAFT, Hirschau, Germany). The casting temperature of the experimental alloys was set at 755 °C ± 5 °C due to the higher Ti addition and poorer solubility of the master alloy AlTi5B1. The chemical composition of the Ti-treated alloys is shown in Table 2. From Table 2, it can be concluded that the actual Ti content in the melt was lower due to the incomplete melting of the AlTi5B1 master alloy.

2.1. Evaluation of the Susceptibility to Tearing

After casting, samples of the experimental alloys were subjected to a quantitative and qualitative evaluation of the susceptibility to tearing. For the experimental work, 4 experimental alloys were chosen. For each experimental alloy, 5 samples were cast to investigate the susceptibility of the alloy to hot tearing.
The design of the mold (Figure 1a), which was part of the overall measuring device (Figure 1b), was designed to support the formation of tears in the casting during solidification. The gating system transported the melt into the cavity of the mold, which was formed by five arms of different lengths. The ends of 4 of the 5 arms were made with spherical ends so that these ends would support the occurrence of inhibited shrinkage and cause a higher susceptibility to tear formation. These 4 arms were used for quantitative evaluation of the tearing susceptibility. The quantitative analysis was performed visually and then verified by capillary control for tear detection. The fifth test arm was used for qualitative evaluation. The fifth (longest) arm was provided with an anchor screw at the end so that its part was overcast with melt. The other end of the screw was attached to the S9M load cell (HBM, Darmstadt, Germany), whose purpose was to record the tensile forces acting in the casting during solidification. A thermocouple was placed in the mold cavity and connected to the NATIONAL INSTRUMENTS Hi-Speed USB Carrier NI USB 9162 measurement card (National Instrument, Austin, TX, USA), which processed the data. The purpose of the thermocouple was to record the temperature as the alloys solidified. From this, the temperature and time of tear initiation could be determined in the event of a tear occurrence in the casting.
For each experimental alloy, a quantitative evaluation was performed on 20 arms of different lengths with a spherical ends and 5 qualitative evaluations of the longest arm. No automatic images inspection software was used in the quantitative and qualitative evaluation.

2.1.1. Quantitative Evaluation

In the first step, the susceptibility of the experimental alloys to tearing was evaluated quantitatively. The quantitative evaluation was based on the determination of the so-called Hot Tearing Index (HTI). For the quantitative assessment of the tearing susceptibility, 4 arms of different lengths were determined. The arms were ended with a spherical end to ensure the occurrence of shrinkage. The Hot Tear Index value was determined using 2 different methods (HTI1 and HTI2) and their combination (HTI3).
The first method for obtaining the Hot Tearing Index value in this study was described by Wu [23]. The HTI1 value is determined on the basis of the number of tears (NT), the type and size of the tear (“weighting factor”—WF) and the number of evaluated arms (NA) by the following Equation (1).
HTI1 = (NT × ΣWF)/NA
The second method for determining the value of HTI2 was described in the work of Akhyar and Song [8,13]. In this case, the value of HTI2 was determined depending on the type and size of the tear (WF), the location of the tear on the test arm (tear position coefficient—TPC), and the length of the test arm (arm length coefficient—ALC) as follows:
HTI2 = ΣWF × TPC × ALC
The coefficient of the tear type and size WF in the evaluation of HTI1 and HTI2 indicated the severity of the tear. Li presented a categorization system that divides the WF coefficient into 4 categories according to the severity of the tear (Figure 2) [24].
The relevant value of the WF coefficient was determined by visual inspection of the tears on 4 test arms. Capillary control was performed to confirm the nature of the tears. The capillary control was performed according to EN ISO 3452-1 [25]. A test system marked IICe according to EN ISO 3452-1 was used for the capillary control. A Diffu-Therm capillary system consisting of a BRE cleaner, BDR red penetrant and BEA white developer (Diffu-Therm, Herten, Germany) was used for the color capillary control. The penetration time was 60 min. It was evaluated in two steps. The first evaluation was performed immediately after the developer dried and the second after the development time of 15 min. The test was performed at a temperature of 22 °C.
Table 3 shows the values of the arm length coefficient (ALC) and the tear position coefficient (TPC) for determining the HTI2 value.
By combining the calculations for determining the values of HTI1 and HTI2, a third method for calculating the value of HTI3 was obtained according to Relationship (3). The HTI3 value provides a more comprehensive view of the alloy’s susceptibility to tearing.
HTI3 = (NT × Σ(WF × ALC × TPC))/NA
The resulting Hot Tear Index values (HTI1, HTI2, and HTI3) can be characterized as a measure of the hot-tearing susceptibility (HTS) (Table 4).

2.1.2. Qualitative Evaluation

After the quantitative evaluation, a qualitative evaluation of the alloy’s susceptibility to tearing was performed. The qualitative evaluation was performed by analyzing the force distribution on the anchor bolt in the longest test arm as a function of the time. Based on this analysis, it was possible to detect a violation of the integrity of the test arm and possibly the formation of tears. The rate of the increase curve was obtained by taking the first derivative of the force curve. This curve was used to determine the temperature at which tearing occurs.

2.2. Thermal Analysis

Thermal analysis was performed to evaluate the solidification process and determine the solidification interval of the experimental alloys. The experimental alloys were cast in a cylindrical mold with a diameter of 35 mm and a height of 50 mm. A type K thermocouple was placed at the bottom of the mold to record the solidification temperature of the melt. The temperature curve of the melt solidification was recorded on a measurement card, which allowed us to convert the acquired dataset into an analog notation. The analog notation was then converted into a numerical and graphical notation using LabView 2 Hz software (version 18.5, National Instrument, Austin, TX, USA). The cooling curve of the experimental alloys as a function of the time was generated using LabView 2 Hz software version 18.5. The crystallization temperature of each phase is difficult to identify from the cooling curve. For this reason, a first derivation of the cooling curve was performed. The crystallization temperature of each phase was characterized by a change in the course of the curve on the first derivative of the cooling curve.

2.3. Microstructural Analysis

The hot tear profiles and fracture surfaces of the experimental alloys were evaluated with a NEOPHOT 2 optimal microscope (OM) (Carl Zeiss AG, Jena, Germany) and TESCAN LMU II scanning electron microscope with a BRUKER EDX analyzer (Bruker Quantax EDX analyzer, Bunker, Kalkar, Germany). The experimental samples were prepared by wet hand-grinding, polishing on polishing wheels that had been moistened with alcohol and impregnated with diamond paste, and then polishing with a Struers Laboforce-3 automatic polisher (Struers, Prague, Czech Republic). The experimental samples were etched with 0.5% HF. Microstructural evaluation was performed on the test arms where immediate tear-off occurred.
For the microstructural evaluation, the effect of the graded Ti addition on the grain refinement was assessed metallographically by evaluating the dendrite arm spacing (DAS) index using Quick Photo Industrial 3.1 software (PROMICRA, Prague, Czech Republic). The DAS index determines the spacing of the secondary axes of the dendrites, and the smaller the DAS index value, the finer the structure and the smaller the segregation spacing. The DAS index value is obtained as the ratio of the dendrite length (L) to the number of dendrite secondary axes (n):
DAS = L/(n − 1)

3. Results

3.1. Thermal Analysis

The thermal analysis record of the experimental alloys is shown in Figure 3.
The solidification of the reference alloy AlSi5Cu2Mg starts with the crystallization of the primary α-(Al) phase. The beginning of the crystallization of the primary α-(Al) phase was recorded at a temperature of 610 °C. Zhang investigated the solidification process of Al-Si alloys [26]. Zhang stated [26] that the solidification of Ti-refined Al-Si alloys can be described as follows:
L → Ti-rich phases + primary phase α-(Al) + eutectic Al-Si
Based on the above, it can be seen from the thermal analysis record that the solidification of the Ti-refined AlSi5Cu2Mg alloy starts with the crystallization of the Ti-based phases [18]. The crystallization of the Ti phases takes place before the solidification of the primary phase α-(Al) in the temperature range of 638 to 645 °C (it is also the liquidus temperature).
The crystallization temperatures of the respective phases obtained from the thermal analysis are processed in Table 5. From the obtained results, it was shown that the grain refining of the AlSi5Cu2Mg-Ti alloy resulted in an increase in the liquidus temperature (TL). This fact is consistent with Farahany’s studies [27]. The liquidus temperature of the alloy without Ti addition was 610 °C. The highest increase in the liquidus temperature was recorded for the alloy with 0.3 wt.% Ti addition (Ti-3). In this case, the liquid temperature increased by 35 °C compared to the reference alloy. By subsequently increasing the wt.% of Ti, an increase in the TL was recorded. An increase in the liquid temperature due to the addition of Ti is accompanied by an increase in the solidification interval of the AlSi5Cu2Mg alloy, which may lead to an increase in the susceptibility of the experimental alloy to tearing [28].
The solidification of the Ti-added alloys continued with the crystallization of the Al-Si eutectic phase at a temperature of about 553 °C. Compared to the reference alloy, the crystallization temperature of the Al-Si eutectic phase decreased by 14 °C. Solidification ends at the solidus temperature (TS). The TS decreased with the addition of Ti compared to the reference alloy, while the largest decrease of 36 °C was recorded for the Ti-2 alloy. From the above, it can be concluded that the addition of Ti to the AlSi5Cu2Mg alloy affects the crystallization temperature of the respective phases [26].

3.2. Quantitative Evaluation of Susceptibility to Tearing

A visual inspection of the test arms was carried out prior to the quantitative evaluation to characterize the severity of the tears produced. The visual inspection was subsequently verified by capillary control (Figure 4). The detected tears were marked with red circles in Figure 4.

3.2.1. Determination of HTI1 Value

The values of the HTI1 index for the reference alloy and Ti alloy are shown in Figure 5. Based on the evaluation of the number, type and size of the tears on the respective test arms, the index for the reference alloy was calculated to be HTI1 = 1.91. From the point of view of evaluating the degree of susceptibility to HTS, the alloy was included in the group of alloys with moderate susceptibility to tearing. The alloys with the addition of Ti showed a slight decrease in the value of the HTI1 index, while its value decreased with increasing wt.% Ti. By adding 0.10 wt.% Ti, the value of the evaluated index decreased by 14% compared to the reference alloy. The value of the HTI1 index = 1.48 was obtained with the Ti-2 alloy, where the decrease with respect to the reference was 22%. The lowest value was obtained with the alloy containing 0.3 wt.% Ti. The value of the HTI1 index of the Ti-3 alloy decreased from 1.91 to 1.41 (a decrease of 26%). Despite the slight decrease in the HTI1, the alloys with Ti addition were classified as alloys with moderate susceptibility to tearing.

3.2.2. Determination of HTI2 Value

The values of the HTI2 index for the reference alloy and the Ti-added alloy are shown in Figure 6. The second method for determining the HTI2 took into account, in addition to the type and size of the tear, the location of the tear on the test arm and the length of the test arm. From this point of view, the HTI2 values obtained were significantly higher than those of the HTI1 index. The reference alloy reached the value of HTI2 index = 3.45 and was included among the alloys with a very high susceptibility to tearing. The alloys with the addition of Ti up to 0.2 wt.% Ti recorded a decrease in the value of the HTI2 index. Compared to the reference alloy, the Ti-1 and Ti-2 alloys showed a decrease in the HTI2 index of 6.66% and 12.71%, respectively. Ti-1 and Ti-2 were classified as high-susceptibility alloys. The HTI2 index determined for the Ti-3 alloys was close to the HTI2 index determined for the reference alloy. The Ti-3 alloy, as well as the alloy without Ti addition, was classified as an alloy with a very high susceptibility to tearing.

3.2.3. Determination of HTI3 Value

The value of the HTI3 index was obtained by combining the relationships used to calculate HTI1 and HTI2 (Figure 7). The alloy without Ti addition reached the value of HTI3 = 3.15 and was classified among the alloys with high susceptibility to tearing. Subsequently, a decrease in the value of the HTI3 index was observed with an increasing wt.% of Ti. With the addition of 0.1 wt.% of Ti, the HTI3 index decreased from 3.15 to 2.92 (a decrease of 7.30%). The Ti-2 alloy showed a 13.97% decrease in HTI3 compared to the reference alloy. In this case, the best results were obtained with the Ti-3 alloy. The value of the HTI3 index decreased by almost 28% from 3.15 to 2.28 with the addition of 0.3 wt.% Ti. The alloys with Ti additions were classified as having a high susceptibility to tearing, as was the reference alloy.

3.3. Qualitative Evaluation of the Susceptibility to Tearing

The results of the qualitative evaluation of the reference alloy are presented in Table 6. The table is separated into two parts. The first part of the table defines the tear initiation conditions: temperature, tear initiation time, maximum force applied to the anchor bolt placed in the test arm during alloy solidification, and the rate of force increase. The second part of the table provides information on the type of damage to the test arm (violation of arm integrity, immediate arm fracture, arm undamaged). For qualitative measurements no. 1 and 5, the occurrence of tears was not detected on the graphical record. Measurement no. 3 showed the formation of a tear. Qualitative evaluation of the reference alloy showed the immediate fracture of the arm in two specimens (#2 and #4).
For the evaluation of the force progression and the rate of force increase, the alloys from Test #1 (no fracture) and Test #2 (immediate fracture of the arm) were used. Figure 8a shows a recorded load curve of the reference alloy from Test #1, in which no tear formation was detected in the test arm. From the graph of the load curve, it can be seen that there are no stresses in the melt prior to the onset of solidification. After the start of solidification, a crystalline structure and deformations occur as the temperature gradually decreases (black curve). An increase in tensile strength was observed on the load curve as the temperature of the alloy decreased. The tensile force increases due to the shrinkage of the casting (blue curve). The green curve, obtained by taking the first derivative of the solidification curve with respect to the time, shows the beginning of the development of the rate of load increase. Figure 8c shows the reference alloy from the first measurement.
Figure 8b shows the load history of the reference alloy from Test #2. In this case, compared to the alloy from Test #1, there was no increase in strength due to the stresses generated during the solidification of the casting. Based on the force curve, it could be assumed that the test arm was immediately torn off along the entire cross-section during solidification. The alloy from Test #2 with the arm immediately torn off is shown in Figure 8d.
The results of the qualitative evaluation of the alloys with Ti addition are shown in Table 7, Table 8 and Table 9. The tables provide a summary of the information on the conditions of tear formation and the type of damage to the test arm. Based on the data presented in Table 6, it can be shown that in no case was the formation of a serious tear, which would lead to the immediate separation of the arm, recorded in the case of the Ti-1 alloy. Out of five qualitative measurements, two alloys (measurements #3 and #4) showed the formation of a tear on the test arm. There was no violation of the integrity of the arm for the alloys from measurements #1, #2, and #3. The results of the qualitative evaluation of the Ti-2 alloy (Table 7) show that in specimens no. 1–3, there was no violation of the integrity of the arm. For the Ti-2 alloys, the formation of a tear was recorded in measurements #4 and #5, but the tear was not severe enough to cause the arm to break immediately. Of the qualitative measurements performed for the Ti-3 alloy (Table 8), only measurement #1 showed tear formation. In measurements no. 2–4, there was no violation of the integrity of the test arm.
From the qualitative measurements performed, the alloys Ti-1 (measurement no. 3), Ti-2 (measurement no. 5) and Ti-3 (measurement no. 2) were selected for further evaluation. In Figure 9b, the Ti-1 alloy from measurement no. 3 is shown. Visual inspection of the alloy from Test #3 revealed a violation of the integrity of the arm. The visual inspection of the test arm was verified by capillary control, which showed no tears in the arm. Based on the graphical record of the load history (Figure 9a) and the results of the visual and capillary control, the presence of an internal tear in the test arm was assumed. On the graph (Figure 9a), red arrows indicate the time of tear initiation and the end of tear propagation. At the end of tear propagation, an increase in the force acting on the anchor bolt to a maximum of 324 N was recorded.
In the Ti-2 alloy, measurement no. 5, the formation of a capillary tear was recorded (Figure 9d). The graphical record of the load is shown in Figure 9b. The area of tear initiation and propagation is marked with red arrows. The tear propagation was stopped after 9 s from the tear initiation. After the end of tear propagation, an increase in the force acting on the anchor bolt was observed. The maximum force achieved on the anchor bolt was 287 N. The force on the anchor bolt, due to the presence of a tear, was significantly lower than that of the alloys from test no. 1–3, where no violation of shoulder integrity was detected.
In Figure 10a, the recorded load curve of the Ti-3 alloy from Test no. 2 is shown, which is characterized by a smooth course of the force curve without tear detection. The maximum value achieved for the force acting on the bolt in the test arm was 974 N. The green curve, obtained by the first derivative of the solidification curve with respect to the time, indicates the beginning of the development of the rate of load increase. Figure 10b shows the Ti-3 alloy from test no. 2.
By comparing the data obtained from the quantitative and qualitative evaluation of the susceptibility to tearing of the experimental alloys, the positive effect of the addition of Ti in reducing the susceptibility to tearing was confirmed. The addition of Ti to the AlSi5Cu2Mg alloy resulted in the grain refinement of the α-(Al) phase, which was one of the main factors influencing the reduction in tearing susceptibility. The same results were presented in the work of [29]. It was based on the evaluation of the influence of the Al-Ti-B inoculant in the AlSi6Cu4 alloy on the susceptibility to tearing, which was significantly reduced after the addition of the inoculant [29]. Another study [30] confirmed the beneficial effect of the AlTi5B1 inoculant in Al alloys in reducing the susceptibility to hot tearing. By using the inoculant, the susceptibility to hot tearing was reduced by changing the grain morphology from columnar to equiaxed and reducing the equiaxed grain size [31]. This improved the overall homogeneity of the structure and allowed better deposition of the melt in the interdendritic spaces, which was also associated with higher resistance to tearing.
Thermal analysis of the experimental alloys showed that the addition of Ti results in an extension of the solidification interval of the AlSi5Cu2Mg alloy (Table 10). The extension of the solidification interval is accompanied by an increase in the susceptibility of the alloy to hot tearing [28]. Despite the extension of the solidification interval of the alloys with Ti addition, these alloys showed a lower susceptibility to hot tearing than the reference alloy. The decisive positive factor in this case was the softening effect of Ti on the primary phase α-(Al), which improved the material’s filling ability in the interdendritic spaces [31].

3.4. Microstructural Evaluation of Hot-Tearing Relief

To evaluate the microstructure of the experimental alloys, samples were selected in which the integrity of the arm was broken during solidification and its immediate separation was observed. By visual inspection, immediate separation of the arm was observed for the alloy without Zr addition, Ti-1 and Ti-3 alloys. Microstructural evaluation was performed in the area where the tear occurred on the arm.
By observing the microstructure of the fracture profile of the reference alloy without Ti addition, it was found that the fracture occurred due to ductile failure along the dendrite boundaries of the α-(Al) phase (Figure 11a,b). One of the initiation sites for tear formation was contractions that occurred in the interdendritic regions of the material (Figure 11a). In addition to ductile failure during tear initiation, areas of cleavage failure were observed on the tear profile (Figure 11c) due to the presence of intermetallic phases based on Cu and Fe.
Microstructural evaluation of the Ti-1 alloy profile showed that the tear initiation mechanism was characterized by ductile failure along the dendrite boundaries of the α-(Al) phase. Locally, a fracture was observed at the boundary between the dendrites of the α-(Al) phase and the Cu-based intermetallic phases (Figure 12a). The fracture surface of the Ti-1 alloy is characterized by brittle fracture due to the presence of cleavage facets. By means of qualitative surface analysis, the cleavage facets were identified as intermetallic phases rich in Cu and Fe (Figure 12b). Figure 12c shows the EDX analysis of the Cu- and Fe-based intermetallic phases, which served as a suitable site for transcrystalline tear propagation.
Evaluation of the microstructure from the tear profile of the Ti-3 alloy showed that the tear mechanism was characterized by ductile failure along the dendrite boundaries of the α-(Al) phase (Figure 13a). The tear profile revealed the presence of hard and simultaneously brittle Cu-based intermetallic phases that could initiate tear propagation. Fractographic evaluation showed the presence of cleavage facets on the fracture surface of the tear, which were identified by surface EDX analysis as intermetallic phases based on Cu, Fe and Ti (Figure 13b). Visual assessment showed that the relief of the fracture surface of the tear was dominated by brittle fracture. SEM analysis of the tear profile at the tear site revealed the presence of small sharp-edged particles. These particles were identified as Ti- and Cu-rich phases by point EDX analysis (Figure 13c). It is believed that the addition of Ti in an excess amount (0.3 wt.%) resulted in the precipitation of a greater number of Ti-rich phases, which tended to precipitate along the grain boundaries of the primary α-(Al) phase.
The metallographic evaluation of the DAS index was performed on the reference alloy and Ti-added alloys using Quick Photo Industrial 3.1 software. The DAS index values of the reference alloy and the Ti-added alloys are shown in Table 11. The values shown are the average of five DAS index measurements.
The reference alloy achieved a DAS index value of 26.8. The addition of 0.1 wt.% Ti resulted in a 21% decrease in the DAS index compared to the reference alloy. The DAS index value of the Ti-0.2 alloy decreased by 24% compared to the reference alloy. The decrease in the DAS index value indicates the achievement of a finer structure due to the grain refinement effect of Ti [32]. A further increase in the wt.% of Ti did not show a significant decrease in the DAS index, which is consistent with studies reporting that increasing Ti above 0.20 wt.% does not result in the significant refinement of the structure [18].

4. Conclusions

This paper focused on evaluating the effect of graded Ti addition on the hot-tearing susceptibility of a hypoeutectic AlSi5Cu2Mg alloy. The Ti addition was deliberately chosen to exceed the manufacturer’s recommendation under the assumption that the hot-tearing susceptibility of the AlSi5Cu2Mg alloy would be reduced due to the grain refinement effect. Based on the results obtained, the following conclusions can be formulated:
  • The solidification interval of the experimental alloys with Ti addition was significantly prolonged compared to the reference alloy. The extension of the solidification interval may increase the susceptibility to hot tearing.
  • Quantitative evaluation of the hot-tearing susceptibility showed that the AlSi5Cu2Mg alloy has a high susceptibility to hot tearing. By evaluating the parameters HTI1, HTI2 and HTI3, it was demonstrated that the susceptibility of the AlSi5Cu2Mg alloy to hot tearing decreases due to the addition of Ti.
  • Qualitatively, a positive decrease in the hot-tearing susceptibility of the AlSi5Cu2Mg alloy was observed with the addition of Ti. The addition of Ti resulted in an effective refinement of the primary α-(Al) phase and a transformation of the columnar grains into equiaxial grains, resulting in better melt-filling ability in the interdendritic spaces. The improved melt-filling ability in the interdendritic spaces resulted in higher tear resistance.
  • Microstructural analysis showed that hot tears tended to propagate along the grain boundaries of the primary α-(Al) phase. The mechanism of tear initiation was characterized by ductile failure and the fracture surface was characterized by brittle fracture due to the presence of Fe-, Cu- and Ti-based intermetallic phases.
  • The evaluation of the DAS index showed that there was a decrease in the value of the DAS index due to the addition of Ti. Based on the decrease in the DAS index, a refining effect of Ti on the primary phase α-(Al) was demonstrated.
The investigation showed a positive effect of the above limit of Ti addition in reducing the susceptibility of the AlSi5Cu2Mg alloy to hot tearing. Despite the positive reduction in the hot-tearing susceptibility, the negative effect of excessive Ti addition on the resulting mechanical and physical properties must also be considered [33]. The AlSi5Cu2Mg alloy used in highly stressed cylinder head castings is characterized by an undesirably high susceptibility to hot tearing. In this regard, it is necessary to further investigate ways to reduce the hot-tearing susceptibility of the AlSi5Cu2Mg alloy.

Author Contributions

Conceptualization, M.M., D.B. and M.S.; methodology, M.M. and M.S.; software, M.M. and M.S.; validation, M.M. and D.B.; formal analysis, M.S.; investigation, M.M. and M.S.; resources, D.B.; data curation, M.M. and M.S.; writing—original draft preparation, M.M. and M.S.; writing—review and editing, M.M., D.B. and M.S.; visualization, M.M. and M.S.; supervision, D.B.; project administration, M.M. and D.B.; funding acquisition, D.B. All authors have read and agreed to the published version of the manuscript.

Funding

This article was created as part of the grant agency projects: VEGA 1/0160/22; VEGA 1/0241/23 and KEGA 029ŽU-4/2023. The authors are thankful for the support.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Conflicts of Interest

The authors declare no conflicts of interest.

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Metals 14 00703 g001
Figure 1. (a) Scheme of the mold; and (b) measuring device: 1—mold, 2—anchoring screw, 3—gripping mechanism, 4—load cell, 5—data processing, 6—inlet, 7—thermocouple, 8—refractory glass.
Figure 1. (a) Scheme of the mold; and (b) measuring device: 1—mold, 2—anchoring screw, 3—gripping mechanism, 4—load cell, 5—data processing, 6—inlet, 7—thermocouple, 8—refractory glass.
Metals 14 00703 g001
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Figure 2. Categorization according to the severity of the tear.
Figure 2. Categorization according to the severity of the tear.
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Metals 14 00703 g003
Figure 3. Thermal analysis of the experimental alloys.
Figure 3. Thermal analysis of the experimental alloys.
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Metals 14 00703 g004
Figure 4. Capillary control of experimental samples: (a) reference alloy, and (b) Ti-3 alloy.
Figure 4. Capillary control of experimental samples: (a) reference alloy, and (b) Ti-3 alloy.
Metals 14 00703 g004
Metals 14 00703 g005
Figure 5. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI1.
Figure 5. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI1.
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Metals 14 00703 g006
Figure 6. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI2.
Figure 6. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI2.
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Metals 14 00703 g007
Figure 7. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI3.
Figure 7. Dependence of the hot-tearing susceptibility of the experimental alloys according to the HTI3.
Metals 14 00703 g007
Metals 14 00703 g008aMetals 14 00703 g008b
Figure 8. Curve of the load force, temperature, and load force ratio: (a) ref. alloy (Test #1), and (b) ref. alloy (Test #2); the sample from the respective measurement: (c) ref. alloy (Test #1), and (d) ref. alloy (Test #2).
Figure 8. Curve of the load force, temperature, and load force ratio: (a) ref. alloy (Test #1), and (b) ref. alloy (Test #2); the sample from the respective measurement: (c) ref. alloy (Test #1), and (d) ref. alloy (Test #2).
Metals 14 00703 g008aMetals 14 00703 g008b
Metals 14 00703 g009aMetals 14 00703 g009b
Figure 9. Curve of the load force, temperature, and load force ratio: (a) Ti-1 alloy (Test #3), and (b) Ti-2 alloy (Test #2); the sample from the respective measurement: (c) Ti-1 alloy (Test #3), and (d) Ti-2 alloy (Test #2).
Figure 9. Curve of the load force, temperature, and load force ratio: (a) Ti-1 alloy (Test #3), and (b) Ti-2 alloy (Test #2); the sample from the respective measurement: (c) Ti-1 alloy (Test #3), and (d) Ti-2 alloy (Test #2).
Metals 14 00703 g009aMetals 14 00703 g009b
Metals 14 00703 g010aMetals 14 00703 g010b
Figure 10. (a) Curve of the load force, temperature, and load force ratio of Ti-3 alloy (Test #2); and (b) the sample from the respective measurement of Ti-3 alloy (Test #2).
Figure 10. (a) Curve of the load force, temperature, and load force ratio of Ti-3 alloy (Test #2); and (b) the sample from the respective measurement of Ti-3 alloy (Test #2).
Metals 14 00703 g010aMetals 14 00703 g010b
Metals 14 00703 g011
Figure 11. Evaluation of hot tear formation and propagation for the reference alloy, OM, SEM: (a) hot tear profile; (b) fracture surface; and (c) EDX analysis of intermetallic phase based on Cu and Fe in the hot tear profile.
Figure 11. Evaluation of hot tear formation and propagation for the reference alloy, OM, SEM: (a) hot tear profile; (b) fracture surface; and (c) EDX analysis of intermetallic phase based on Cu and Fe in the hot tear profile.
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Metals 14 00703 g012
Figure 12. Evaluation of hot tear formation and propagation for the Ti-1 alloy, OM, SEM: (a) hot tear profile; (b) fracture surface; and (c) EDX analysis of intermetallic phase based on Cu and Fe in the hot tear profile.
Figure 12. Evaluation of hot tear formation and propagation for the Ti-1 alloy, OM, SEM: (a) hot tear profile; (b) fracture surface; and (c) EDX analysis of intermetallic phase based on Cu and Fe in the hot tear profile.
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Metals 14 00703 g013aMetals 14 00703 g013b
Figure 13. Evaluation of hot tear formation and propagation for the Ti-3 alloy, OM, SEM: (a) hot tear profile, (b) fracture surface, and (c) EDX analysis of intermetallic phase based on Cu and Ti in the hot tear profile.
Figure 13. Evaluation of hot tear formation and propagation for the Ti-3 alloy, OM, SEM: (a) hot tear profile, (b) fracture surface, and (c) EDX analysis of intermetallic phase based on Cu and Ti in the hot tear profile.
Metals 14 00703 g013aMetals 14 00703 g013b
Table 1. Chemical composition of the reference alloy [wt.%].
Table 1. Chemical composition of the reference alloy [wt.%].
SiCuMgFeZrSrTiMoAl
5.491.920.290.190.00090.010.0130.006Bal.
Table 2. Chemical composition of the experimental alloys with the addition of Ti [wt.%].
Table 2. Chemical composition of the experimental alloys with the addition of Ti [wt.%].
SiCuMgTiSrFeMoAl
Ti-15.811.840.270.090.0070.1820.0059Bal.
Ti-25.721.880.230.170.0050.1530.0061Bal.
Ti-35.621.840.220.250.0060.1620.0058Bal.
Table 3. Evaluation system for the ALC and TPC for HTI2.
Table 3. Evaluation system for the ALC and TPC for HTI2.
Arm Length CoefficientALCTear Position CoefficientTPC
Arm 1 (64.5 mm)1Lower part of arm1
Arm 2 (104.5 mm)2Middle part of the arm3
Arm 3 (124.5 mm)3Upper part of arm2
Arm 4 (184.5 mm)4
Table 4. Hot-tearing susceptibility intervals reprinted from Ref. [23].
Table 4. Hot-tearing susceptibility intervals reprinted from Ref. [23].
<0.50.5–1.251.25–2.252.25–3.5>3.5
HTS MinimalLowModerateHighVery high
Table 5. Crystallization temperature of the experimental alloys [°C].
Table 5. Crystallization temperature of the experimental alloys [°C].
Ti-Phasesα-PhaseEutectic Al + SiTS
Ref. alloy-610567517
Ti-1638625554483
Ti-2640623551481
Ti-3645622553486
Table 6. Qualitative evaluation of the reference alloy.
Table 6. Qualitative evaluation of the reference alloy.
No. Hot Tear InitiationEnd of Hot Tear Propagation
Temperature (°C)Time (s)Load (N)Load Force Ratio (N/s)Temperature (°C)Time (s)Type of End Hot Tear Propagation
1.No hot tearmax. 1260max. 7.9No hot tear
2.Immediate arm separationImmediate arm separation
3.47422max. 655max. 4.744033Increase of load
4.Immediate arm separationImmediate arm separation
5.No hot tearmax. 1328max. 12.3No hot tear
Table 7. Qualitative evaluation of the Ti-1 alloy.
Table 7. Qualitative evaluation of the Ti-1 alloy.
No. Hot Tear InitiationEnd Hot Tear Propagation
Temperature (°C)Time (s)Load (N)Load Force Ratio (N/s)Temperature (°C)Time (s)Type of End Hot Tear Propagation
1.No hot tearmax. 1075max. 8.1No hot tear
2.No hot tearmax. 1035max. 7.3No hot tear
3.42548max. 324max. 2.839752Increase of load
4.41928max. 655max. 6.139236Increase of load
5.No hot tearmax. 1145max. 9.5No hot tear
Table 8. Qualitative evaluation of the Ti-2 alloy.
Table 8. Qualitative evaluation of the Ti-2 alloy.
No. Hot Tear InitiationEnd Hot Tear Propagation
Temperature (°C)Time (s)Load (N)Load Force Ratio (N/s)Temperature (°C)Time (s)Type of End Hot Tear Propagation
1.No hot tearmax. 1018max. 7.8No hot tear
2.No hot tearmax. 1168max. 9.7No hot tear
3.No hot tear 2max. 920max. 9.7No hot tear
4.54916max. 253max. 2.551325Increase of load
5.43040max. 1328max. 12.339249Increase of load
Table 9. Qualitative evaluation of the Ti-3 alloy.
Table 9. Qualitative evaluation of the Ti-3 alloy.
No. Hot Tear InitiationEnd Hot Tear Propagation
Temperature (°C)Time (s)Load (N)Load Force Ratio (N/s)Temperature (°C)Time (s)Type of End Hot Tear Propagation
1.51620max. 683max. 4.748726Increase of load
2.No hot tearmax. 974max. 9.5No hot tear
3.No hot tearmax. 1123max. 7.5No hot tear
4.No hot tearmax. 1151max. 10.6No hot tear
5.No hot tearmax. 1145max. 9.5No hot tear
Table 10. The solidification interval of the experimental alloys [°C].
Table 10. The solidification interval of the experimental alloys [°C].
Ref. AlloyTi-1Ti-2Ti-3
610–517638–483640–481645–486
Table 11. Dendrite arm spacing of the experimental alloys (µm).
Table 11. Dendrite arm spacing of the experimental alloys (µm).
Ref. AlloyTi-1Ti-2Ti-3
26.821.120.319.9
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Matejka, M.; Bolibruchová, D.; Sýkorová, M. Effect of Ti Addition on the Hot-Tearing Susceptibility of the AlSi5Cu2Mg Alloy. Metals 2024, 14, 703. https://doi.org/10.3390/met14060703

AMA Style

Matejka M, Bolibruchová D, Sýkorová M. Effect of Ti Addition on the Hot-Tearing Susceptibility of the AlSi5Cu2Mg Alloy. Metals. 2024; 14(6):703. https://doi.org/10.3390/met14060703

Chicago/Turabian Style

Matejka, Marek, Dana Bolibruchová, and Martina Sýkorová. 2024. "Effect of Ti Addition on the Hot-Tearing Susceptibility of the AlSi5Cu2Mg Alloy" Metals 14, no. 6: 703. https://doi.org/10.3390/met14060703

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