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Article

Characterisation of a New Generation of AlMgZr and AlMgSc Filler Materials for Welding Metal–Ceramic Composites

by
Jan Wysocki
1,
Marek Staude
2,*,
Andrzej Trytek
3 and
Mariusz Sosnowski
2
1
Faculty of Marine Engineering, Maritime University of Szczecin, 70-500 Szczecin, Poland
2
Faculty of Mechatronics and Electrical Engineering, Maritime University of Szczecin, 70-500 Szczecin, Poland
3
Faculty of Mechanics and Technology, Rzeszow University of Technology, 37-450 Stalowa Wola, Poland
*
Author to whom correspondence should be addressed.
Appl. Sci. 2024, 14(10), 4177; https://doi.org/10.3390/app14104177
Submission received: 22 April 2024 / Revised: 10 May 2024 / Accepted: 11 May 2024 / Published: 15 May 2024
(This article belongs to the Section Additive Manufacturing Technologies)
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Abstract

:
When manufacturing the welded joints of components made of metal–ceramic composites of the Al-Si/SiC type, we encounter significant difficulties. This is related to the presence of a ceramic phase in the aluminium alloy matrix. The interaction between the molten metal matrix and the ceramic particles in the weld pool influences a complex of physicochemical phenomena resulting in, among other things, the structure of the welded joint. This is particularly true of the effect of the distribution of ceramic particles and their influence on the crystallisation process in the weld pool. An important issue is the influence of the reinforcing particles on the susceptibility of the aluminium matrix to both hot and cold cracking. The scope of the research included the development of the chemical composition of an additive material for the TIG welding of aluminium–ceramic composites. This material was made in the form of so-called sticks, cast from alloys containing elements such as magnesium, scandium or zirconium in addition to aluminium. The appropriate composition of the mass content of the individual components was intended to change the crystallisation mode of the weld pool and to obtain strengthening precipitates. The most favourable structure was obtained in the case of a modification of the AlMg5 alloy by the addition of scandium. Minor dispersions of Al3Sc became the nucleation pads of fine grains, which improved the mechanical properties of the alloy. Also, in the case of the addition of zirconium, the crystallisation shifted from dendritic to fine-grain growth. In this paper, the basic strength properties of the developed materials were tested and the most favourable chemical compositions of the filling materials were selected.

1. Introduction

Metal matrix composites are increasingly used as construction materials due to their special mechanical properties [1,2,3]. Cast aluminium alloy-based composites [4,5,6] reinforced with ceramic material in the form of long fibres [7,8], short fibres [9,10] and particles [11,12] account for a significant share of this application. In addition to the many advantages conditioning their use, they are characterised by certain disadvantages, which primarily include the difficulty of joining structural elements using welding methods [13,14,15]. The mechanical properties and structural structure of joints of aluminium alloy matrix composites in friction welding [16,17] and friction stir welding [18,19,20] have been widely described in the literature. During welding with the above-mentioned methods, material melting and re-solidification do not occur, which minimises the surface reactions at the reinforcing phase–metal matrix interface. Despite its undoubted advantages, friction welding has limitations in terms of the shape of the parts to be joined.
One of the basic welding techniques is the TIG (Tungsten Inert Gas) method, which is relatively inexpensive and allows the joining of parts of varying shapes [21,22,23]. The application of the TIG method for joining aluminium composites reinforced with ceramic particles involves melting the metal matrix and re-solidification of the liquid weld pool. The manner of crystallisation of the matrix alloy and the influence of the crystallisation front on the distribution of the reinforcing phase, wettability and the product of surface reactions at the liquid metal–ceramic phase interface have a significant impact on the mechanical properties of the weld. Typically, welded microstructures are characterised by a dendritic structure with an uneven distribution of the reinforcing phase in the interdendritic spaces [24,25,26,27,28]. In the case of metal–ceramic composites, the factor of the physical and chemical interactions that occur between the liquid metal and the ceramic reinforcing phase must additionally be considered [29,30,31]. The durability of the bond between the liquid and solid phase depends on the surface tension of the liquid and the contact angle, which determines wettability. For aluminium–silicon ceramic systems, the contact angle exceeds 90° and, in most cases, even 120°. Hence, it follows that the work of adhesion in such systems has a low value, which is detrimental to the mechanical properties of the composites [32,33]. One possibility for improving the metal–ceramic interface is to introduce chemical elements into the liquid metal that improve the wettability of the reinforcing phase by reducing the surface tension, lowering the surface energy of the liquid at the liquid–solid interface or forming chemical compounds that improve wettability [34,35]. Many of the aluminium alloys exhibit a tendency for hot cracking during the welding process. The applied heat energy and the associated temperature distribution in the welded joint area results in internal stresses and microstructural changes in the re-solidified weld and zones with partial melting of the material [36]. Internal stresses leading to open hot cracks result from two consecutive processes: the expansion of the metal under the influence of the supplied heat energy and the significant contraction of the metal during crystallization [37]. The second type of damage occurring in the heat-affected zone is the so-called liquid cracking, created by the presence of a film of liquid metal with a low melting point at the grain boundary. As stated by Huang and Kou [38,39], liquid cracking in the HAZ depends on the stability and solubility limit in the α-phase of the liquid-initiating particles of the precipitates formed during the fabrication process. Sub-melts appear at the phase boundaries, after which they solidify in the form of eutectics. The phase segregation created in this way contributes to a reduction in the strength properties of the weld in the heat-affected zones. The degree of segregation and thus the susceptibility to hot cracking is decisively influenced by the difference in the amount of solid fraction in the weld pool and the zone of partial melting during solidification [40]. Cao and Kou [41], based on temperature–solid phase quantity diagrams, studied welded joints of cast Al-Si aluminium alloy components with Al-Mg and Al-Si additive material. They confirmed that hot cracking in cast alloys, as in heat-treatable alloys, depends on the course of the bond crystallisation and the zone of partial melting (the difference in the amount of solid phase).
One of the most effective methods for preventing hot cracking is to achieve a change in structure from dendritic to fine-grained equiaxial during weld crystallisation. This can be achieved by adding certain alloying elements. As stated by Wang et al. [42], vanadium, introduced into pure aluminium, promoted a change from a peritectic to an equiaxial alloy crystallisation mode. The resulting Al10V separations provided an effective barrier to grain growth. In another paper, Wang et al. [43] showed that aluminium alloys containing zirconium in their composition were also characterised by a fine-grained structure with Al3Zr strengthening separations. Joy [44] noted that zirconium not only affects the mode of crystallisation, but also reduces grain growth during the recrystallisation of AlMg alloys. Zakharov [45] and Belov [46] reported an extremely beneficial effect from scandium on the mechanical properties of aluminium alloys by strengthening them with Al3Sc precipitates. The effect of precipitation strengthening can be enhanced by the simultaneous addition of both zirconium and scandium to the aluminium alloy [47]. Czerwiński F. [48] presented the possibility of introducing scandium and caesium simultaneously into alloys. The introduction of caesium made it possible to reduce the proportion of scandium and lower the manufacturing cost while maintaining the fine-grained structure of the aluminium alloy.
Another element that strongly affects the joint structure of composite materials is the high electric arc energy, which can lead to the disintegration of the ceramic reinforcement. and the formation of undesirable brittle reaction products at the reinforcement–metal matrix interface. For this reason, the welding process parameters such as shielding gas output, current intensity or welding speed, and in particular the way in which the filler material formed from the mixture of the filler metal and the metal matrix crystallises and the strengthening precipitates formed in the process, are extremely important for the mechanical properties of the welded joint.
Tests described in the literature to join cast aluminium alloy-based composites by TIG using the filling material AlMg (ER5356) or Al-Si (ER4347), which are used to join unreinforced aluminium alloys, were characterised by a number of welding discontinuities. Given the positive effects of alloying elements such as zirconium and scandium on the mechanical properties of aluminium alloys through precipitation strengthening, it seems reasonable to attempt to modify the commercial filling material with a suitable alloying element composition.

2. Study Material and Methodology

2.1. Fabrication of Additional Materials for AlSi/SiC(p) Bonding by Casting

In order to produce filling materials with a specific composition of chemical composition that could be used in the TIG welding of silumin-based composite materials, a split mould of heat-resistant steel was made (Figure 1). One part of the mould was a flat plate. The other was a plate, in which a filler channel and a casting channel for “sticks” (additional material) of the intended chemical composition were milled. The mould’s interlocking surfaces were milled and polished in order to achieve a proper seal. The two parts of the mould were connected using two M12 screws. Prior to the casting process, the mould was coated with a slurry made from a mixture of zinc white with the addition of a small amount of graphite powder, dissolved in an aqueous glass solution.
Additive materials were produced by melting AlMg5, AlZr15, AlSc2 and Mg feedstock in an induction furnace and direct casting. The feedstock material was placed in a graphite crucible, which was then placed in an induction furnace and melted at 1800 °C +/− 10 °C held for 20 min, with mechanical stirring of the molten metal every 5 min. The resulting liquid melt was poured into a metal mould, preheated to 500 °C. After cooling, the mould was unmoulded and the resulting alloys were additionally subjected to a homogenisation annealing at 500 °C for 12 h in a WK-106-C type resistance furnace in an ambient atmosphere.
From the casting with the shape shown in Figure 2, the section marked A was used for further impact and tensile tests, while the section marked B was used for hardness and microhardness tests. The section marked C (which can be used as an additional material for TIG bonding tests of cast AlSi/SiC(p) composites) was subjected to analysis of the chemical composition and structural structure of the castings made.

2.2. Research Methodology

Specimens were cut from the resulting castings and the samples were prepared by grinding, polishing and surface etching with Keller’s reagent. Structural and chemical analysis of the prepared samples was carried out within 15 min of completion of the digestion process. This was to minimise the oxidation process of the aluminium alloys and its possible impact on the reliability of the results. The microstructure was analysed using light microscopy Neophot 2 with a magnification of 500 times. Further examination of the microstructure and chemical composition was carried out using an XL30LaB6 microscope with a DX-EDAX X-ray analyser.
Hardness and microhardness tests were carried out using the Vickers method on a Hecker HPO 250 Brinell–Vicker hardness tester according to the standard PN-EN ISO 6507-1:2018-05 [49], with a hardness of HV5 at a load of 48.35 N (HV5) in 10 s and microhardness on a Shimadzu microhardness testing machine at an applied load of 0.98 N (HV0.1) in 10 s.
The tensile test was performed on a Galdabini Quasar 600 material testing machine according to Polish Standard PN-EN ISO 6892-1:2016-09 [50] using non-standard specimens of type A1 (Figure 3).
An impact test was carried out using the Charpie method on a Labor Tech CHK 50J-D testing machine according to standard PN-EN ISO 148-1:2017-02 [51] on five specimens for each cast alloy with V-notch with 2 mm notch depth.

3. Microstructure of the Filling Materials

3.1. Filling Material Type AlMgZr

Three types of AlMgZr alloys with different weight percentages of zirconium were cast for the study. Analysis of the chemical composition together with the X-ray spectrum of the cast filling materials showed that the amount of this element was, respectively, 0.35 wt.% (Figure 4a), 0.67 wt.% (Figure 4b) and 0.93 wt.%. (Figure 4c). The magnesium content varied within 5 wt.% in all castings made.
The observation and metallographic analysis of each of the aforementioned alloys was then carried out. Figure 4 shows example images of the microstructure of AlMg5Zr type castings with different zirconium contents. As can be seen in Figure 4a, the addition of this element at a level of 0.35 wt.% slightly induced a fragmentation effect on the dendritic grains. As the percentage of zirconium increased to 0.67 wt.%, a significant comminution of the casting grains could be observed (Figure 5b) and, at 0.93 wt.%, the ratification effect of the grains diminished, causing them to grow slightly again.
In addition, iron precipitates in the characteristic “Chinese script” form were observed in the AlMg5Zr0.67 casting, as well as the formation of a triple eutectic α + Mg2Si + Si with significant intensification for AlMg5Zr0.93 (Figure 6).
Agglomerates of “rod-shaped” (Al,Si)3Zr precipitates with a size of 30–50 μm and chemical composition as in Figure 7 were visible in all castings.
In the AlMgZr castings, small spherical Al3Zr precipitates ranging from 1 μm to 3 μm in size could also be observed, unevenly distributed throughout the casting at grain boundaries and in small amounts in the interdendritic spaces (Figure 8).
It is known that in binary alloys, zirconium remains in the form of a supersaturated solution upon rapid cooling of the casting. Furthermore, the decomposition of the presaturated solid solution manifests itself through the formation of a metastable L12 Al3Zr phase characterised by negligible compatibility with the aluminium matrix. Further annealing at temperatures in the range 400–500 °C leads to transformation of the equilibrium DO23 Al3Zr phase [52]. In multicomponent alloys, the strong Al-Zr bonding of the solid solution is replaced by weaker Mg-Zr reducing the solubility of zirconium. This effect is amplified when silicon is present in the alloy [53]. This may explain the agglomeration of rod-shaped (Al,Si)3Zr precipitates as a product of undissolved zirconium in the aluminium matrix. On the other hand, the location of the spherical Al3Zr precipitates in the spaces of the primary dendritic grains indicates that the zirconium content of the crystallising liquid was sufficient for their nucleation. The number, type of precipitates and their size were probably related to the temperature and annealing time of the homogenising casting. To achieve maximum precipitation hardening by forming high-density Al3Zr phase dispersoids, a two-step annealing would have to be carried out [54]. Given the scientific use of the alloy as a possible additional AlMgZr material, which involves its re-melting, two-step annealing was not carried out.

3.2. Filling Material Type AlMgSc

Based on the two-phase Al-Sc equilibrium diagram defining the solubility limit of scandium in aluminium [55] and for the ternary AlMgSc alloy [56], three types of castings were made with scandium contents in the sub-eutectic, eutectic and super-eutectic ranges.
An analysis of the chemical composition of the resulting alloys is shown in Figure 9. The scandium contents were, respectively, 0.30 wt.% (Figure 9a), 0.61 wt.%. (Figure 9b) and 0.84 wt.% (Figure 9c). As in the additional materials of the AlMgZr type, a similar magnesium content of 5 wt.% was retained here, too.
The microstructure of AlMgSc-type alloys is shown in Figure 10. In the cast AlMg5Sc0.3 type alloy with scandium contents in the sub-eutectic range (Figure 10a), the grains were irregular with a tendency towards growth with few Al3Sc separations. Large separations of intermetallic phases are also evident.
In the case of AlMg5Sc0.61 alloys, where the scandium content was in the peri-eutectic range, the microstructure showed an equiaxial grain structure with phase separations at the grain boundaries. An increased amount of Al3Sc precipitates was also observed (Figure 11), which became shims for further crystallisation of the casting.
With a further increase in the proportion of scandium in the AlMg5Sc0.84-type casting, a tendency towards an increase in the amount of precipitates of the super-eutectic Al3Sc phase and a further, albeit minor, fragmentation of the equiaxial grains became apparent. At the same time, in the AlMg5SC0.61 and AlMg5Sc0.84 castings, a certain amount of small Al3Sc precipitates at the grain boundaries, which were displaced during the crystallisation process, as well as iron precipitates, was observed (Figure 12).
Structural changes in the fabricated AlMgSc additives of 0.61 wt.% and 0.84 wt.%, where the amount of fine Al3Sc precipitates after grain boundaries probably depended on the cooling rate of the casting mould. The amount of fine precipitates across the grain boundaries of the equiaxial grains increased with keeping the casting in the mould in the elevated temperature range and extending the annealing period. The annealing period itself did not affect the recrystallisation of the castings as might be expected, since scandium at concentrations above 0.6 wt.% added to AlMg alloys acts as a strong modifier and significantly increases the recrystallisation temperature due to the high dispersity of secondary Al3Sc precipitates distributed throughout the casting volume [57].

4. Testing of Mechanical Properties of Filling Materials

4.1. Testing of Hardness and Microhardness

The hardness and microhardness test results are shown in Figure 13. As can be seen, the addition of zirconium or scandium to the AlMg5 alloy significantly increased both parameters. Additionally, it was noted that the hardness was highest at 0.35 wt.% and dropped for a content of 0.67 wt.% and slightly increased for 0.93 wt.% but did not achieve the same value as the AlMg5Zr0.35 alloy. The increase in hardness of AlMgZr filling materials was mainly due to the presence of fine Al3Zr precipitate located on the grain boundaries.
The addition of scandium to the AlMg5 base material also significantly increased the hardness parameter. A further significant increase in microhardness was observed for the casting up to a scandium content of 0.61 wt.%. A further increase in the scandium content to 0.84 wt.% resulted in only a slight increase in this parameter. However, the value of the standard deviation of the measurements taken does not make it clear whether this value has changed or whether it is the result of measurement errors.

4.2. Static Tensile Strength Test

The tensile test results obtained are shown in Figure 14.
A cast alloy of the AlMg5Sc0.3 type, where the zirconium content was too low for significant grain size reduction, reached a tensile strength of 135 MPa, with 1.5% ductility. This effect was to be expected, as it is typical of castings with large grain sizes with a eutectic phase at their boundaries. A clear increase in tensile strength was obtained for castings with a scandium content higher than 0.6 wt.%; there was a 10% improvement in tensile strength, with a concomitant increase in ductility. However, for scandium contents of 0.84 wt.%, the tensile strength increase parameter was only increased by a further 2.5 per cent. It was also observed that the ductility of the AlMg5Sc0.84 alloy type deteriorated slightly, which may be due to the intensification of Al3Sc precipitation.
The cast AlMgZr-type alloys showed a similar trend. Tensile strength increased with the increasing zirconium content in the alloy. The proportion of zirconium in the alloy at 0.35 wt.% did not significantly affect the tensile strength parameter. This is probably due to zirconium segregation during the crystallisation process. A further increase in the zirconium proportion resulted in an increase in the strength and ductility of the alloy, which could be explained by an increase in the proportion of Al3Zr-segregated particles in the total volume of the materials tested, as well as favouring grain fragmentation.

4.3. Impact Test

The test results obtained and presented in Table 1 show that, although a slight increase in impact strength was observed in the AlMgZr castings with increased zirconium mass content, the statement from the obtained values is not conclusive. The recorded fracture work for the individual samples had a rather high standard deviation.
In the AlMgSc castings, the individual samples had a much lower standard deviation and the results appear more conclusive. The impact test showed that the work of fracture increased with the increasing scandium proportion to a value of 0.61 wt.%. Further increases in the mass proportion of scandium in the alloy only slightly increased the impact strength of the casting.

5. Discussion

It is known that alternating current TIG welding of aluminium alloys is accompanied by so-called cathodic cleaning to break up the surface oxide layer. The high arc energy also leads to the partial disintegration of the ceramic reinforcing phase of the fused parent metal and contributes to the formation of harmful products that impede the wettability of the ceramics through the metal matrix [58]. Thus, the addition of alloying elements to improve weldability, reducing the amount of eutectic phase and changeing the crystallisation mode from dendritic to fine-grained appears to be an effective way to improve weld quality in composite materials. This can minimise the process of ceramic particles being pushed out by the growth of the dendritic crystallisation front, thereby changing the mechanical properties of the welded joint [13]. The modification of the AlMg5 alloy by the addition of scandium or zirconium changed its structural structure. Dispersion of the intermetallic phase Al3Sc in AlMgSc alloys or Al3Zr in AlMgZr alloys forms the nucleation pads of fine grains. The fine-grained structure and secretion strengthening favourably shape the course of weld pool crystallisation, resulting in good mechanical weld properties close to those of the parent material [59,60]. Tests carried out on AlMgZr and AlMgSc alloys have shown that:
  • The addition of zirconium to the AlMg5 alloy reduced dendritic grain growth and, through Al3Zr precipitates, improved its mechanical properties;
  • The zirconium in the alloy formed two types of precipitates in the casting. One was in the form of “sticks” and the other in the form of dobby precipitates unevenly distributed in the eutectic near grain boundaries, which had the effect of increasing hardness and microhardness;
  • The addition of Sc to AlMg5 alloys changed the structure of the casting to a fine-grained structure. This effect was most pronounced with the content of this element corresponding to the eutectic value for the Al-Sc two-phase alloy;
  • The Al3Sc precipitates distributed in the central part of the grains and in the form of very small precipitates along the grain boundaries visibly reduced the amount of the eutectic phase. This alloy structure increased the tensile strength;
  • Considering the results obtained from the tests carried out, it can be seen that the addition of scandium improved the mechanical properties of the AlMg alloys to a greater extent than had the zirconia modification. Furthermore, it is known that the addition of scandium significantly increases corrosion resistance and improves weldability [44,45,46,47]. The only aspect that may cause a limitation in the wider use of this element as a modifier of AlMg alloy-based filling material is its high price. Therefore, zirconium, although offering slightly inferior mechanical properties, is less expensive and can also be used as an alloying additive to improve the mechanical properties of TIG welded joints.

Author Contributions

Conceptualisation, J.W.; methodology, J.W.; software, M.S. (Marek Staude); validation, J.W., A.T., M.S. (Marek Staude) and M.S. (Mariusz Sosnowski); formal analysis, J.W.; investigation, A.T.; resources, J.W. and M.S. (Marek Staude); data curation, M.S. (Marek Staude); writing—original draft preparation, M.S. (Marek Staude); writing—review and editing, M.S. (Mariusz Sosnowski); visualisation, M.S. (Mariusz Sosnowski); supervision, A.T.; project administration, M.S. (Marek Staude); funding acquisition, J.W. All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The original contributions presented in the study are included in the article, further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

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Applsci 14 04177 g001
Figure 1. Macroscopic view of the casting mould structure for filler material production.
Figure 1. Macroscopic view of the casting mould structure for filler material production.
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Figure 2. Macroscopic view of one of the received castings.
Figure 2. Macroscopic view of one of the received castings.
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Figure 3. Dimensions of the custom sample used for the tensile strength test of the resulting castings.
Figure 3. Dimensions of the custom sample used for the tensile strength test of the resulting castings.
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Figure 4. X-ray microanalysis of the chemical composition of the filling materials cast: (a) AlMg5Zr0.35; (b) AlMg5Zr0.67; (c) AlMg5Zr0.93.
Figure 4. X-ray microanalysis of the chemical composition of the filling materials cast: (a) AlMg5Zr0.35; (b) AlMg5Zr0.67; (c) AlMg5Zr0.93.
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Figure 5. Microstructure of cast filling materials of AlMgZr type: (a) AlMg5Zr0.35; (b) AlMg5Zr0.67; (c) AlMg5Zr0.93.
Figure 5. Microstructure of cast filling materials of AlMgZr type: (a) AlMg5Zr0.35; (b) AlMg5Zr0.67; (c) AlMg5Zr0.93.
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Figure 6. X-ray microanalysis of the chemical composition of the precipitate phases of the AlMg5Zr0.93 casting.
Figure 6. X-ray microanalysis of the chemical composition of the precipitate phases of the AlMg5Zr0.93 casting.
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Figure 7. Cluster of club-shaped precipitates in an AlMgZr0.84 casting.
Figure 7. Cluster of club-shaped precipitates in an AlMgZr0.84 casting.
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Figure 8. Al3Zr precipitates in an AlMgZr0.67 casting.
Figure 8. Al3Zr precipitates in an AlMgZr0.67 casting.
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Figure 9. X-ray microanalysis of the chemical composition of the filling materials cast: (a) AlMg5Sc0.30; (b) AlMg5Sc0.61; (c) AlMg5Sc0.84.
Figure 9. X-ray microanalysis of the chemical composition of the filling materials cast: (a) AlMg5Sc0.30; (b) AlMg5Sc0.61; (c) AlMg5Sc0.84.
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Figure 10. Microstructure of the cast Al-Mg-Sc additive materials: (a) AlMg5Sc0.3; (b) AlMg5Sc0.61; (c) AlMg5Sc0.84.
Figure 10. Microstructure of the cast Al-Mg-Sc additive materials: (a) AlMg5Sc0.3; (b) AlMg5Sc0.61; (c) AlMg5Sc0.84.
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Figure 11. Al3Sc precipitates in an AlMgSc0.63 alloy.
Figure 11. Al3Sc precipitates in an AlMgSc0.63 alloy.
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Figure 12. Precipitates of fine Al3Sc phases at grain boundaries and the presence of iron precipitates in the AlMgSc0.84 casting.
Figure 12. Precipitates of fine Al3Sc phases at grain boundaries and the presence of iron precipitates in the AlMgSc0.84 casting.
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Figure 13. Graphical representation of the results of HV5 hardness and HV0.1 microhardness measurement of the produced AlMgZr and AlMgSc binders and the standard AlMg5 binder. (*)—value provided by the manufacturer.
Figure 13. Graphical representation of the results of HV5 hardness and HV0.1 microhardness measurement of the produced AlMgZr and AlMgSc binders and the standard AlMg5 binder. (*)—value provided by the manufacturer.
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Figure 14. Tensile strengths of cast AlMgZr and AlMgSc alloys.
Figure 14. Tensile strengths of cast AlMgZr and AlMgSc alloys.
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Table 1. Impact test results for as-cast filler materials.
Table 1. Impact test results for as-cast filler materials.
Filling MaterialInitial
Energy [J]		Specimen
Length [mm]		Specimen
Width [mm]		V Notch
Depth [mm]		Impact
Energy [J]		Impact Strength
[J/cm2]		Standard Deviation
AlMg5Zr0.3549.03325101024.82977512.074440.222583
AlMg5Zr0.674.78074311.951850.367749
AlMg5Zr0.934.80525912.013150.346717
AlMg5Sc0.3049.03325101024.78074211.951860.200849
AlMg5Sc0.614.90332512.258310.190333
AlMg5Sc0.844.92784212.319610.174729
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MDPI and ACS Style

Wysocki, J.; Staude, M.; Trytek, A.; Sosnowski, M. Characterisation of a New Generation of AlMgZr and AlMgSc Filler Materials for Welding Metal–Ceramic Composites. Appl. Sci. 2024, 14, 4177. https://doi.org/10.3390/app14104177

AMA Style

Wysocki J, Staude M, Trytek A, Sosnowski M. Characterisation of a New Generation of AlMgZr and AlMgSc Filler Materials for Welding Metal–Ceramic Composites. Applied Sciences. 2024; 14(10):4177. https://doi.org/10.3390/app14104177

Chicago/Turabian Style

Wysocki, Jan, Marek Staude, Andrzej Trytek, and Mariusz Sosnowski. 2024. "Characterisation of a New Generation of AlMgZr and AlMgSc Filler Materials for Welding Metal–Ceramic Composites" Applied Sciences 14, no. 10: 4177. https://doi.org/10.3390/app14104177

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