<figure><img src="image_650.png" /><figcaption>Fig. 9 Appearance of extruded product with chip billets of various densities</figcaption></figure>

## Conclusions

The following conclusions are drawn based on the hot extrusion, extrusion simulation, and surface observation of AA6063 chip material using a solid die and a pseudo porthole die.

1) Chip materials can be solidified to a density close to that of AA6063 by pressing at 200 °C and 100 t.

2) In the case of extrusion using a pseudo porthole die, the extrusion force was approximately 1.5 times higher than that using a solid die.

3) A higher strain was generated in the material using the pseudo porthole die than in the material using the solid die.

4) In the hot extrusion process of chip billets, the use of a pseudo porthole die suppressed the defects on the surface of the product, which occurs when a solid die is used.

5) When the density of chip billets approached that of AA6063, extruded materials without surface defects were obtained.

6) For a low density chip billets, product surface defects occurred until the middle stage of the extrusion; however, no product surface defects occurred in the latter half of the extrusion.

Future research will focus on evaluating the mechanical properties of the extruded chip billets for each density and internal porosity defects such as blister tests. In addition, the port shape of the pseudo porthole die will be varied to determine the relationship between the amount of shear strain applied to the interior and the billet density.

## References

[1] F. Czerwinski, Current trends in automotive lightweighting strategies and materials. Materials, 14(21), (2021), 6631. https://doi.org/10.3390/ma14216631

[2] R. Modaresi, D. B. Müller, The role of automobiles for the future of aluminum recycling, Environ. Sci. Technol., 46(16), (2012), 8587-8594. https://doi.org/10.1021/es300648w

[3] K. Martinsen, S. Gulbrandsen-Dahl, Use of post-consumer scrap in aluminium wrought alloy structural components for the transportation sector, Procedia CIRP, 29, (2015), 686-691. https://doi.org/10.1016/j.procir.2015.02.072
[4] A. Wagiman, M. S. Mustapa, R. Asmawi, S. Shamsudin, M. A. Lajis, Y. Mutoh, A review on direct hot extrusion technique in recycling of aluminium chips, Int. J. Adv. Manuf. Technol., 106, (2020), 641-653. https://doi.org/10.1007/s00170-019-04629-7

[5] M. Samuel, A new technique for recycling aluminium scrap, J. Mater. Process. Technol., 135(1), (2003), 117-124. https://doi.org/10.1016/S0924-0136(02)01133-0

[6] S. Capuzzi, G. Timelli, Preparation and melting of scrap in aluminum recycling: A review, Metals, 8(4), (2018), 249. https://doi.org/10.3390/met8040249

[7] Y. Tabata, Aluminum VISION 2050: Outline and implications, J. Jpn. Inst. Light. Met, 72, (2022), 549-555. https://doi.org/10.2464/jilm.72.549

[8] A. Schulze, O. Hering, A. E. Tekkaya, Production and Subsequent Forming of Chip-Based Aluminium Sheets Without Remelting, Int. J. of Precis. Eng. and Manuf.-Green Tech., 9(4), (2022), 1035-1048. https://doi.org/10.1007/s40684-021-00395-8

[9] S. N. Ab Rahim, M. A. Lajis, S. Ariffin, A review on recycling aluminum chips by hot extrusion process, Procedia CIRP, 26, (2015), 761-766. https://doi.org/10.1016/j.procir.2015.01.013

[10] G. Buffa, D. Campanella, M. Adnan, U. La Commare, G. Ingarao, L. Fratini, Enhancing industrial applicability of aluminum alloys chips recycling through Friction Stir Extrusion: Thin wires production process, Manuf. Lett., 35, (2023), 371-378. https://doi.org/10.1016/j.mfglet.2023.08.131

[11] R. Chiba, M. Yoshimura, Solid-state recycling of aluminium alloy swarf into c-channel by hot extrusion, J. Manuf. Process., 17, (2015), 1-8. https://doi.org/10.1016/j.jmapro.2014.10.002

[12] T. Funazuka, N. Takatsuji, T. Tsuchiya, S. Oda, Pick-up defect mechanism in hot extrusion of Al-Mg-Si series alloy. J. Jpn. Inst. Light. Met., 70(9), (2020), 415-421. https://doi.org/10.2464/jilm.70.415

[13] A. E. Tekkaya, M. Schikorra, D. Becker, D. Biermann, N. Hammer, K. Pantke, Hot profile extrusion of AA-6060 aluminum chips, J. Mater. Process. Technol., 209(7), (2009), 3343-3350. https://doi.org/10.1016/j.jmatprotec.2008.07.047

[14] V. Güley, A. Güzel, A. Jäger, N. Ben Khalifa, A. E. Tekkaya, W. Z. Misiolek, Effect of die design on the welding quality during solid state recycling of AA6060 chips by hot extrusion, Mater. Sci. Eng. A, 574, (2013), 163-175. https://doi.org/10.1016/j.msea.2013.03.010

[15] M. Haase, A. E. Tekkaya, Recycling of aluminum chips by hot extrusion with subsequent cold extrusion, Procedia Eng., 81, (2014), 652-657. https://doi.org/10.1016/j.proeng.2014.10.055

[16] T. Funazuka, K. Dohda, M. Kinoshita, S. Suzuki, Phenomena of tool adhesion at elevated temperature in V-groove friction test of AA7075, In Materials Research Proceedings: Material Forming: ESAFORM 2023, 28, (2023), 419-428. https://doi.org/10.21741/9781644902479-46
[4] A. Wagiman, M. S. Mustapa, R. Asmawi, S. Shamsudin, M. A. Lajis, Y. Mutoh, A review on direct hot extrusion technique in recycling of aluminium chips, Int. J. Adv. Manuf. Technol., 106, (2020), 641-653. https://doi.org/10.1007/s00170-019-04629-7

[5] M. Samuel, A new technique for recycling aluminium scrap, J. Mater. Process. Technol., 135(1), (2003), 117-124. https://doi.org/10.1016/S0924-0136(02)01133-0

[6] S. Capuzzi, G. Timelli, Preparation and melting of scrap in aluminum recycling: A review, Metals, 8(4), (2018), 249. https://doi.org/10.3390/met8040249

[7] Y. Tabata, Aluminum VISION 2050: Outline and implications, J. Jpn. Inst. Light. Met, 72, (2022), 549-555. https://doi.org/10.2464/jilm.72.549

[8] A. Schulze, O. Hering, A. E. Tekkaya, Production and Subsequent Forming of Chip-Based Aluminium Sheets Without Remelting, Int. J. of Precis. Eng. and Manuf.-Green Tech., 9(4), (2022), 1035-1048. https://doi.org/10.1007/s40684-021-00395-8

[9] S. N. Ab Rahim, M. A. Lajis, S. Ariffin, A review on recycling aluminum chips by hot extrusion process, Procedia CIRP, 26, (2015), 761-766. https://doi.org/10.1016/j.procir.2015.01.013

[10] G. Buffa, D. Campanella, M. Adnan, U. La Commare, G. Ingarao, L. Fratini, Enhancing industrial applicability of aluminum alloys chips recycling through Friction Stir Extrusion: Thin wires production process, Manuf. Lett., 35, (2023), 371-378. https://doi.org/10.1016/j.mfglet.2023.08.131

[11] R. Chiba, M. Yoshimura, Solid-state recycling of aluminium alloy swarf into c-channel by hot extrusion, J. Manuf. Process., 17, (2015), 1-8. https://doi.org/10.1016/j.jmapro.2014.10.002

[12] T. Funazuka, N. Takatsuji, T. Tsuchiya, S. Oda, Pick-up defect mechanism in hot extrusion of Al-Mg-Si series alloy. J. Jpn. Inst. Light. Met., 70(9), (2020), 415-421. https://doi.org/10.2464/jilm.70.415

[13] A. E. Tekkaya, M. Schikorra, D. Becker, D. Biermann, N. Hammer, K. Pantke, Hot profile extrusion of AA-6060 aluminum chips, J. Mater. Process. Technol., 209(7), (2009), 3343-3350. https://doi.org/10.1016/j.jmatprotec.2008.07.047

[14] V. Güley, A. Güzel, A. Jäger, N. Ben Khalifa, A. E. Tekkaya, W. Z. Misiolek, Effect of die design on the welding quality during solid state recycling of AA6060 chips by hot extrusion, Mater. Sci. Eng. A, 574, (2013), 163-175. https://doi.org/10.1016/j.msea.2013.03.010

[15] M. Haase, A. E. Tekkaya, Recycling of aluminum chips by hot extrusion with subsequent cold extrusion, Procedia Eng., 81, (2014), 652-657. https://doi.org/10.1016/j.proeng.2014.10.055

[16] T. Funazuka, K. Dohda, M. Kinoshita, S. Suzuki, Phenomena of tool adhesion at elevated temperature in V-groove friction test of AA7075, In Materials Research Proceedings: Material Forming: ESAFORM 2023, 28, (2023), 419-428. https://doi.org/10.21741/9781644902479-46
# Layer adhesion and critical strain of HFCVD diamond coatings on WC-Co substrate

LECHNER Stefan¹,ᵃ*, BARON Sarah²,ᵇ, HÖFER Markus²,ᶜ, STEIN Christian²,ᵈ, MÜLLER Sören¹,ᵉ

¹Extrusion Research & Development Center, Technische Universität Berlin, Gustav-Meyer-Allee 25, 13355 Berlin, Germany

²Fraunhofer Institute for Surface Engineering and Thin Films IST, Riedenkamp 2, 38108 Braunschweig, Germany

ᵃs.lechner@tu-berlin.de, ᵇsarah.baron@ist.fraunhofer.de, ᶜmarkus.hoefer@ist.fraunhofer.de, ᵈchristian.stein@ist.fraunhofer.de, ᶜsoeren.mueller@tu-berlin.de

**Keywords:** Compression Tests, Cobalt Content, Mechanical Properties, Interface Stability

**Abstract.** The extrusion of products with high demands on dimensional accuracy and surface quality, such as aluminum power bonding wires or bioabsorbable stents made of magnesium, renders high requirements on the extrusion die. With the application of HFCVD (Hot Filament Chemical Vapor Deposition) diamond coatings on cobalt-based tungsten carbide alloys, a die material with great hot strength, high surface hardness and low chemical adhesion could be provided. To identify WC-Co alloys with suitable cobalt content both for use as extrusion die and for coating with CVD diamond, investigations on ductility and hot strength as well as layer adhesion and interface stability were carried out. The mechanical properties of four WC-Co alloys containing 6%, 9%, 10% and 12% cobalt were determined by means of hot compression tests. Layer adhesion was evaluated by a standardized sandblasting test. Subsequently, hot compression tests were carried out with coated specimens to determine the critical deformation of the applied CVD diamond coatings.

**Introduction**

In aluminum extrusion, the production of delicate profiles such as micro-multiport profiles (MMP) for heat exchangers is a well-established process. In the case of products with similar dimensions but high dimensional accuracy requirements, such as aluminum power bonding wires or bio-absorbable stents made of magnesium, only semi-finished products have been produced by extrusion up to now. The reason for this is the high demands placed on the mechanical and tribological properties of the extrusion dies. High extrusion ratios and resulting high extrusion pressures require elevated hot strength of the die materials. Furthermore, very low dimensional tolerances (<10 µm) and optimum surface quality of the profiles necessitate considerable reduction of die wear.

The billet temperature in the extrusion of aluminum and magnesium is usually between 300-500 °C and is therefore within the optimum application range of conventional tool steels [1]. However, the high specific extrusion pressures of over 1200 MPa when extruding profiles with dimensions in the micrometer range clearly exceed their high-temperature strength. Elastic and plastic deformations are induced in the tools, which have a negative effect on the dimensional accuracy of the extruded products. Even high-performance materials such as nickel- and cobalt-based alloys, which have excellent high-temperature strength, only achieve strengths of up to 1100 MPa in this temperature range and therefore do not meet these extreme requirements [2]. Hard metals, on the other hand, exhibit compressive strengths of up to 8000 MPa. They consist of

<figure><img src="image_640.png" /></figure>
# Layer adhesion and critical strain of HFCVD diamond coatings on WC-Co substrate

LECHNER Stefan¹,ᵃ*, BARON Sarah²,ᵇ, HÖFER Markus²,ᶜ, STEIN Christian²,ᵈ, MÜLLER Sören¹,ᵉ

¹Extrusion Research & Development Center, Technische Universität Berlin, Gustav-Meyer-Allee 25, 13355 Berlin, Germany

²Fraunhofer Institute for Surface Engineering and Thin Films IST, Riedenkamp 2, 38108 Braunschweig, Germany

ᵃs.lechner@tu-berlin.de, ᵇsarah.baron@ist.fraunhofer.de, ᶜmarkus.hoefer@ist.fraunhofer.de, ᵈchristian.stein@ist.fraunhofer.de, ᶜsoeren.mueller@tu-berlin.de

**Keywords:** Compression Tests, Cobalt Content, Mechanical Properties, Interface Stability

**Abstract.** The extrusion of products with high demands on dimensional accuracy and surface quality, such as aluminum power bonding wires or bioabsorbable stents made of magnesium, renders high requirements on the extrusion die. With the application of HFCVD (Hot Filament Chemical Vapor Deposition) diamond coatings on cobalt-based tungsten carbide alloys, a die material with great hot strength, high surface hardness and low chemical adhesion could be provided. To identify WC-Co alloys with suitable cobalt content both for use as extrusion die and for coating with CVD diamond, investigations on ductility and hot strength as well as layer adhesion and interface stability were carried out. The mechanical properties of four WC-Co alloys containing 6%, 9%, 10% and 12% cobalt were determined by means of hot compression tests. Layer adhesion was evaluated by a standardized sandblasting test. Subsequently, hot compression tests were carried out with coated specimens to determine the critical deformation of the applied CVD diamond coatings.

**Introduction**

In aluminum extrusion, the production of delicate profiles such as micro-multiport profiles (MMP) for heat exchangers is a well-established process. In the case of products with similar dimensions but high dimensional accuracy requirements, such as aluminum power bonding wires or bio-absorbable stents made of magnesium, only semi-finished products have been produced by extrusion up to now. The reason for this is the high demands placed on the mechanical and tribological properties of the extrusion dies. High extrusion ratios and resulting high extrusion pressures require elevated hot strength of the die materials. Furthermore, very low dimensional tolerances (<10 µm) and optimum surface quality of the profiles necessitate considerable reduction of die wear.

The billet temperature in the extrusion of aluminum and magnesium is usually between 300-500 °C and is therefore within the optimum application range of conventional tool steels [1]. However, the high specific extrusion pressures of over 1200 MPa when extruding profiles with dimensions in the micrometer range clearly exceed their high-temperature strength. Elastic and plastic deformations are induced in the tools, which have a negative effect on the dimensional accuracy of the extruded products. Even high-performance materials such as nickel- and cobalt-based alloys, which have excellent high-temperature strength, only achieve strengths of up to 1100 MPa in this temperature range and therefore do not meet these extreme requirements [2]. Hard metals, on the other hand, exhibit compressive strengths of up to 8000 MPa. They consist of

<figure><img src="image_651.png" /></figure>
a hard material, usually carbide-based, which is embedded in a metal matrix. The metal matrix significantly increases the toughness of the composite material but reduces the hardness [3]. For extrusion dies, primarily inserts made of cobalt-based tungsten carbide hard meals (WC-Co) are used. Hard metals also have the advantage of being very wear-resistant [4], which makes them meet both requirements for the planned application [5].

The wear occurring on the tool surfaces is composed of abrasive and adhesive wear. Abrasive wear primarily influences the shape and dimensional accuracy of the profile cross-section, while adhesive wear in the extrusion channel of the die reduces the surface quality of the profile [1]. In the case of iron-based tool materials, the high adhesion tendency of aluminum and magnesium to iron causes local welds to form between billet and die material during extrusion. In the bearing channel, these cause scoring ("die line defect") on the profile produced [6]. The adhesion tendency between the material pairings is therefore a decisive factor for the surface quality of the extruded profile.

In addition to the application of wear-resistant materials, surface treatment of extrusion dies offers an effective way of keeping product quality constant over many extrusion cycles. Conventional surface modifications such as nitriding or CVD coating with TiC, TiN or Al2O3 delay the formation of adhesion, but do not completely prevent them [1,7]. Diamond coatings are unique in terms of their combination of extreme hardness, wear resistance and chemical inertness. Especially in contact with aluminum, diamond behaves largely inert and shows the lowest wetting compared to other CVD coatings such as TiN or Al2O3 [8].

In this study, the suitability of WC-Co with HFCVD (Hot Filament CVD) diamond coating for use as extrusion dies was investigated. Ductility and hot strength of four different WC-Co alloys were characterized by means of hot compression tests. To evaluate layer adhesion and interface stability of the diamond coatings on those alloys, a standardized sandblasting test was carried out. A further set of hot compression tests was performed to determine the critical strain of different diamond coatings as a function of coating morphology and thickness, as well as substrate pretreatment.

## Experimental Procedure

The suitability of WC-Co for coating with CVD diamond and for use as an extrusion die poses opposing requirements for the cobalt content. Higher cobalt content favors ductility [3] and therefore the use as die but impairs application of CVD diamond [9]. To identify WC-Co alloys with suitable cobalt content both for use as extrusion die and for coating with CVD diamond, investigations on ductility and hot strength as well as layer adhesion and interface stability were carried out (see Figure 1).

<figure><img src="image_641.png" /><figcaption>Figure 1: Schematic illustration of the experimental procedure: a) hot compression tests of WC-Co alloys, b) HFCVD coating of specimen, c) sandblasting test and d) hot compression tests of coated WC-Co specimen.</figcaption></figure>

Hot compression tests on WC-Co alloys. Concerning the mechanical properties, hot compression tests were performed on four different alloys containing 6 %, 9 %, 10 % and 12 % cobalt (EMT100, EMT609, EMT210, EMT612). The average grain size for the WC-particles was 0.8 µm (EMT100, EMT210) and 0.5 µm (EMT609, EMT612). Cylinders with a diameter of Ø5 mm and a length of 7.5 mm were tested in a thermal-mechanical physical simulation system
a hard material, usually carbide-based, which is embedded in a metal matrix. The metal matrix significantly increases the toughness of the composite material but reduces the hardness [3]. For extrusion dies, primarily inserts made of cobalt-based tungsten carbide hard meals (WC-Co) are used. Hard metals also have the advantage of being very wear-resistant [4], which makes them meet both requirements for the planned application [5].

The wear occurring on the tool surfaces is composed of abrasive and adhesive wear. Abrasive wear primarily influences the shape and dimensional accuracy of the profile cross-section, while adhesive wear in the extrusion channel of the die reduces the surface quality of the profile [1]. In the case of iron-based tool materials, the high adhesion tendency of aluminum and magnesium to iron causes local welds to form between billet and die material during extrusion. In the bearing channel, these cause scoring ("die line defect") on the profile produced [6]. The adhesion tendency between the material pairings is therefore a decisive factor for the surface quality of the extruded profile.

In addition to the application of wear-resistant materials, surface treatment of extrusion dies offers an effective way of keeping product quality constant over many extrusion cycles. Conventional surface modifications such as nitriding or CVD coating with TiC, TiN or Al2O3 delay the formation of adhesion, but do not completely prevent them [1,7]. Diamond coatings are unique in terms of their combination of extreme hardness, wear resistance and chemical inertness. Especially in contact with aluminum, diamond behaves largely inert and shows the lowest wetting compared to other CVD coatings such as TiN or Al2O3 [8].

In this study, the suitability of WC-Co with HFCVD (Hot Filament CVD) diamond coating for use as extrusion dies was investigated. Ductility and hot strength of four different WC-Co alloys were characterized by means of hot compression tests. To evaluate layer adhesion and interface stability of the diamond coatings on those alloys, a standardized sandblasting test was carried out. A further set of hot compression tests was performed to determine the critical strain of different diamond coatings as a function of coating morphology and thickness, as well as substrate pretreatment.

## Experimental Procedure

The suitability of WC-Co for coating with CVD diamond and for use as an extrusion die poses opposing requirements for the cobalt content. Higher cobalt content favors ductility [3] and therefore the use as die but impairs application of CVD diamond [9]. To identify WC-Co alloys with suitable cobalt content both for use as extrusion die and for coating with CVD diamond, investigations on ductility and hot strength as well as layer adhesion and interface stability were carried out (see Figure 1).

<figure><img src="image_652.png" /><figcaption>Figure 1: Schematic illustration of the experimental procedure: a) hot compression tests of WC-Co alloys, b) HFCVD coating of specimen, c) sandblasting test and d) hot compression tests of coated WC-Co specimen.</figcaption></figure>

Hot compression tests on WC-Co alloys. Concerning the mechanical properties, hot compression tests were performed on four different alloys containing 6 %, 9 %, 10 % and 12 % cobalt (EMT100, EMT609, EMT210, EMT612). The average grain size for the WC-particles was 0.8 µm (EMT100, EMT210) and 0.5 µm (EMT609, EMT612). Cylinders with a diameter of Ø5 mm and a length of 7.5 mm were tested in a thermal-mechanical physical simulation system
Gleeble 3800. Temperatures of 300 °C, 400 °C and 500 °C were applied to the specimens by means of resistance heating. Compression was performed at a strain rate of 0.0001 s⁻¹ until the specimens fractured. For each temperature and WC-Co alloy 3-5 repetitions were carried out.

CVD diamond coating of test specimen. The CVD diamond coating was carried out in a laboratory scale Hot Filament Chemical Vapor Deposition (HFCVD) system at Fraunhofer IST. The coating system enables diamond coatings on areas up to 200 mm x 240 mm. In advance to diamond deposition the cemented carbide samples were cleaned. The various WC-Co alloys were then chemically etched in order to remove the cobalt binder at a depth of a few micrometers in the surface peripheral zone, which interferes with the adhesion of the diamond coating. After ultrasonic treatment in a nano-diamond suspension for enabling diamond seeding of the surfaces, the test specimens (cylinders and flat test bodies) were put into the deposition reactor and placed on a flat substrate holder. The deposition was carried out at filament temperatures of 2150°C in a gas atmosphere consisting of hydrogen and different amounts of methane. For the deposition of microcrystalline diamond films, a pressure of 20 mbar and a CH₄ concentration of 1.6 % was used. In contrast, the nanocrystalline diamond films were deposited at 5 mbar and a CH₄ concentration of 4.0 %. During the deposition process diamond growth took place at substrate temperatures of approx. 800°C, and closed and adhering diamond films were formed. Film thickness was controlled by adjustment of the deposition time. The flat samples used for the evaluation of the adhesion by applying a standardized sandblasting test were coated with nanocrystalline diamond with a film thickness of $d_{film} = 11$ µm. The cylindrical specimens for the hot compression tests were coated with nano- and microcrystalline diamond films, each with thickness values of $d_{film} = 6$ µm and 12 µm, respectively. Since the cylinders were placed with the cylinder surface on the sample holder during diamond deposition, they were only coated on approximately one half, which was directly exposed to the process atmosphere, while the rear segments of the cylinder surfaces remained uncoated.

To be able to distinguish the differently coated samples, they were laser-engraved with individual sample numbers. The planar samples were labelled on the uncoated back sides. For the cylindrical samples, the laser engravement was applied to the front faces.

Evaluation of the adhesion strength and interface stability of deposited CVD diamond films on cemented carbides. The adhesion strength of the deposited diamond films was evaluated on flat test specimens (Ø = 18 mm, h = 5 mm) by applying a standardized sandblasting test. In this test, the diamond-coated surface is bombarded with a very intense beam of SiC particles (grit size F180). The test conditions (pressure, working distance, etc.) are selected in such a way that a good differentiation of diamond coatings with different levels of adhesion can be achieved with test times of maximum 1 to 3 minutes. The time until the first damage to the diamond coating appears is measured and provides a qualitative measure of the coating adhesion and interface stability. Since the test duration is also influenced by the coating thickness, all samples must be coated to the same thickness.

Hot compression tests on CVD-diamond coated WC-Co. A second set of hot compression tests was then carried out with diamond coated cylinders containing 6 % (EMT100) and 9 % (EMT609) cobalt to determine the critical strain, before spalling of the coating occurs. In addition, different modifications of the diamond coating were applied. Morphology (nano/micro-crystalline), coating thickness (6 µm/12 µm) and substrate pretreatment (ground/sandblasted) were varied. The strain rate was again 0.0001 s⁻¹ and temperature was kept constant at 300 °C for all trials. For each combination of substrate material and coating variation, the applied strain was increased iteratively. In the first stage, a specimen was compressed to 50 % of the fracture strain of the substrate material. If there was no visible degradation of the coating or fracture of the specimen, a strain 75 % of the fracture strain was applied to a new specimen with identical coating type and thickness in a second iteration stage. Analogously, the strain was increased to 87.5 % of the
Gleeble 3800. Temperatures of 300 °C, 400 °C and 500 °C were applied to the specimens by means of resistance heating. Compression was performed at a strain rate of 0.0001 s⁻¹ until the specimens fractured. For each temperature and WC-Co alloy 3-5 repetitions were carried out.

CVD diamond coating of test specimen. The CVD diamond coating was carried out in a laboratory scale Hot Filament Chemical Vapor Deposition (HFCVD) system at Fraunhofer IST. The coating system enables diamond coatings on areas up to 200 mm x 240 mm. In advance to diamond deposition the cemented carbide samples were cleaned. The various WC-Co alloys were then chemically etched in order to remove the cobalt binder at a depth of a few micrometers in the surface peripheral zone, which interferes with the adhesion of the diamond coating. After ultrasonic treatment in a nano-diamond suspension for enabling diamond seeding of the surfaces, the test specimens (cylinders and flat test bodies) were put into the deposition reactor and placed on a flat substrate holder. The deposition was carried out at filament temperatures of 2150°C in a gas atmosphere consisting of hydrogen and different amounts of methane. For the deposition of microcrystalline diamond films, a pressure of 20 mbar and a CH₄ concentration of 1.6 % was used. In contrast, the nanocrystalline diamond films were deposited at 5 mbar and a CH₄ concentration of 4.0 %. During the deposition process diamond growth took place at substrate temperatures of approx. 800°C, and closed and adhering diamond films were formed. Film thickness was controlled by adjustment of the deposition time. The flat samples used for the evaluation of the adhesion by applying a standardized sandblasting test were coated with nanocrystalline diamond with a film thickness of $d_{film} = 11$ µm. The cylindrical specimens for the hot compression tests were coated with nano- and microcrystalline diamond films, each with thickness values of $d_{film} = 6$ µm and 12 µm, respectively. Since the cylinders were placed with the cylinder surface on the sample holder during diamond deposition, they were only coated on approximately one half, which was directly exposed to the process atmosphere, while the rear segments of the cylinder surfaces remained uncoated.

To be able to distinguish the differently coated samples, they were laser-engraved with individual sample numbers. The planar samples were labelled on the uncoated back sides. For the cylindrical samples, the laser engravement was applied to the front faces.

Evaluation of the adhesion strength and interface stability of deposited CVD diamond films on cemented carbides. The adhesion strength of the deposited diamond films was evaluated on flat test specimens (Ø = 18 mm, h = 5 mm) by applying a standardized sandblasting test. In this test, the diamond-coated surface is bombarded with a very intense beam of SiC particles (grit size F180). The test conditions (pressure, working distance, etc.) are selected in such a way that a good differentiation of diamond coatings with different levels of adhesion can be achieved with test times of maximum 1 to 3 minutes. The time until the first damage to the diamond coating appears is measured and provides a qualitative measure of the coating adhesion and interface stability. Since the test duration is also influenced by the coating thickness, all samples must be coated to the same thickness.

Hot compression tests on CVD-diamond coated WC-Co. A second set of hot compression tests was then carried out with diamond coated cylinders containing 6 % (EMT100) and 9 % (EMT609) cobalt to determine the critical strain, before spalling of the coating occurs. In addition, different modifications of the diamond coating were applied. Morphology (nano/micro-crystalline), coating thickness (6 µm/12 µm) and substrate pretreatment (ground/sandblasted) were varied. The strain rate was again 0.0001 s⁻¹ and temperature was kept constant at 300 °C for all trials. For each combination of substrate material and coating variation, the applied strain was increased iteratively. In the first stage, a specimen was compressed to 50 % of the fracture strain of the substrate material. If there was no visible degradation of the coating or fracture of the specimen, a strain 75 % of the fracture strain was applied to a new specimen with identical coating type and thickness in a second iteration stage. Analogously, the strain was increased to 87.5 % of the
fracture strain in a third iteration stage. In each iteration stage, the specimens were analyzed by means of optical light microscopy to evaluate the state of the CVD diamond coating.

## Results and Discussion

Ductility and hot strength of WC-Co alloys. The mechanical behavior of WC-Co alloys is highly influenced by cobalt content and grain size [10]. To find a balance of a favorable binder content for both the coating (rigidity) and for use as an extrusion die (elasticity), hot compression tests were carried out on WC-Co alloys with different cobalt contents. The compressive yield strain (true strain) of the tested alloys is shown in Figure 2a as a function of the test temperature. Due to the brittle material behavior of the cemented carbides, the measured values scatter widely [11]. In general, an inverse dependence of the compressive yield strain on temperature can be seen for all materials. The determined strain is reduced with increasing test temperature. The only exception to this is the alloy EMT100, where a lessened strain was measured at 300 °C. On average, all four materials exhibit a similar compressive yield strain of 0.026. A slightly reduced average value was determined for EMT210 (0.023) and a moderately increased average value was found for EMT609 (0.0295). As shown in Figure 2b, the effect of temperature on the compressive yield strength is also inverse. Analogous to the compressive yield strain, an unexpectedly low compressive yield strength was measured for EMT100 at 300 °C. On average, the lowest compressive yield strength was determined for EMT210 (ca. 2780 MPa) and EMT612 (ca. 2890 MPa), which is due to their high binder contents. For EMT100 a slightly elevated average of ca. 3530 MPa was measured and EMT609 revealed the highest compressive yield strength with an average of ca. 4690 MPa. Considering the ultrafine grain of EMT 609 and EMT612, the increased strength of those alloys in comparison with their counterparts with similar cobalt content can be explained [12].

<figure><img src="image_642.png" /><figcaption>Figure 2: a) Compressive yield strain and b) compressive yield strength of the investigated WC-Co alloys.</figcaption></figure>

These results are also in good agreement with the room temperature hardness of the WC-Co alloys [13]. In addition to EMT 100, the alloy EMT609 was selected for further investigations concerning the critical strain of CVD diamond coatings on WC-Co substrate due to its increased compressive yield strain and compressive yield strength in the temperature range of 300-500°C.

Adhesion strength of the deposited CVD diamond films on flat WC-Co samples. A total of 22 hard metal disks with 18 mm diameter were coated with a nanocrystalline diamond layer with thickness of 11 µm. The coated samples were tested by applying the sandblasting test under identical conditions. Figure 2 shows the times to the appearance of the first damage of the diamond coating as a function of the specified Co content of the WC grades EMT100, EMT609, EMT210, and EMT612. If no flaking occurred, the test was terminated after 300 s at the latest. The selected test conditions were extremely abrasive. With any other hard coating such as diamond like carbon
fracture strain in a third iteration stage. In each iteration stage, the specimens were analyzed by means of optical light microscopy to evaluate the state of the CVD diamond coating.

## Results and Discussion

Ductility and hot strength of WC-Co alloys. The mechanical behavior of WC-Co alloys is highly influenced by cobalt content and grain size [10]. To find a balance of a favorable binder content for both the coating (rigidity) and for use as an extrusion die (elasticity), hot compression tests were carried out on WC-Co alloys with different cobalt contents. The compressive yield strain (true strain) of the tested alloys is shown in Figure 2a as a function of the test temperature. Due to the brittle material behavior of the cemented carbides, the measured values scatter widely [11]. In general, an inverse dependence of the compressive yield strain on temperature can be seen for all materials. The determined strain is reduced with increasing test temperature. The only exception to this is the alloy EMT100, where a lessened strain was measured at 300 °C. On average, all four materials exhibit a similar compressive yield strain of 0.026. A slightly reduced average value was determined for EMT210 (0.023) and a moderately increased average value was found for EMT609 (0.0295). As shown in Figure 2b, the effect of temperature on the compressive yield strength is also inverse. Analogous to the compressive yield strain, an unexpectedly low compressive yield strength was measured for EMT100 at 300 °C. On average, the lowest compressive yield strength was determined for EMT210 (ca. 2780 MPa) and EMT612 (ca. 2890 MPa), which is due to their high binder contents. For EMT100 a slightly elevated average of ca. 3530 MPa was measured and EMT609 revealed the highest compressive yield strength with an average of ca. 4690 MPa. Considering the ultrafine grain of EMT 609 and EMT612, the increased strength of those alloys in comparison with their counterparts with similar cobalt content can be explained [12].

<figure><img src="image_653.png" /><figcaption>Figure 2: a) Compressive yield strain and b) compressive yield strength of the investigated WC-Co alloys.</figcaption></figure>

These results are also in good agreement with the room temperature hardness of the WC-Co alloys [13]. In addition to EMT 100, the alloy EMT609 was selected for further investigations concerning the critical strain of CVD diamond coatings on WC-Co substrate due to its increased compressive yield strain and compressive yield strength in the temperature range of 300-500°C.

Adhesion strength of the deposited CVD diamond films on flat WC-Co samples. A total of 22 hard metal disks with 18 mm diameter were coated with a nanocrystalline diamond layer with thickness of 11 µm. The coated samples were tested by applying the sandblasting test under identical conditions. Figure 2 shows the times to the appearance of the first damage of the diamond coating as a function of the specified Co content of the WC grades EMT100, EMT609, EMT210, and EMT612. If no flaking occurred, the test was terminated after 300 s at the latest. The selected test conditions were extremely abrasive. With any other hard coating such as diamond like carbon
(DLC), TiC, TiN, or TiAlN, coating failure occurs within fractions of a second. If the particle beam is directed at uncoated hard metal surfaces, a hole several 100 micrometers in depth is created in just a few seconds. Figure 2 shows that the achievable test times decrease almost exponentially with the cobalt content. Our experiences with sandblasting tests of diamond-coated cutting tools show, that test times of one minute and longer indicate very good to excellent adhesion values. Such performance values were obtained for EMT100 and EMT609 samples with Co contents of 6% and 9%. Both WC-alloys with cobalt contents of 10% and 12% perform significantly worse. Nevertheless, EMT210 with 10% cobalt still mostly reaches durability values of more than 30 seconds, which is sufficient for applications with medium requirements. Even on the EMT612 samples with 12% Co, an adhesive diamond coating was deposited in all cases. The test times achieved are between 6 s and 22 s, which is probably only sufficient for applications with low coating adhesion requirements. Based on the sandblasting test results, we decided to evaluate only EMT100 and EMT609 as substrate materials for extrusion dies and for the hot compression tests shown in the next section.

<figure><img src="image_643.png" /><figcaption>Figure 3: Results of sandblasting tests for nanocrystalline diamond coatings deposited by HFCVD on cemented carbide grades with different cobalt contents. The test duration was limited to a maximum of 300 s.</figcaption></figure>

Critical strain of CVD-diamond-coatings. In order to determine the limiting deformation of the CVD coatings, hot compression tests were carried out using the WC-Co alloys EMT100 and EMT609 as substrate materials. The applied strain was derived from the previous hot compression tests of uncoated samples. Here, an average fracture strain of 0.0333 and 0.0325 was measured for EMT100 and EMT609, respectively. Therefore, the targeted strain values for the three iterations were set to 0.017 (50 % fracture strain), 0.025 (75 % fracture strain) and 0.029 (87,5 % fracture strain) for the three iteration stages. The first substrate material to be tested was EMT100. In the first iteration stage, no defects of the CVD coating were observed. Regardless of the coating configuration (morphology/thickness) or substrate pre-treatment, no visually recognizable damage to the coating occurred. Figure 4a illustrates a specimen, which was tested in the first iteration stage. In the second iteration stage, repeated failure of the substrate material occurred even though the applied strain was below the determined fracture strain. An example is shown in Figure 4b, where the specimen is split in half. Due to the early fracture of the WC-Co substrate, the
(DLC), TiC, TiN, or TiAlN, coating failure occurs within fractions of a second. If the particle beam is directed at uncoated hard metal surfaces, a hole several 100 micrometers in depth is created in just a few seconds. Figure 2 shows that the achievable test times decrease almost exponentially with the cobalt content. Our experiences with sandblasting tests of diamond-coated cutting tools show, that test times of one minute and longer indicate very good to excellent adhesion values. Such performance values were obtained for EMT100 and EMT609 samples with Co contents of 6 % and 9 %. Both WC-alloys with cobalt contents of 10 % and 12 % perform significantly worse. Nevertheless, EMT210 with 10 % cobalt still mostly reaches durability values of more than 30 seconds, which is sufficient for applications with medium requirements. Even on the EMT612 samples with 12 % Co, an adhesive diamond coating was deposited in all cases. The test times achieved are between 6 s and 22 s, which is probably only sufficient for applications with low coating adhesion requirements. Based on the sandblasting test results, we decided to evaluate only EMT100 and EMT609 as substrate materials for extrusion dies and for the hot compression tests shown in the next section.

<figure><img src="image_654.png" /><figcaption>Figure 3: Results of sandblasting tests for nanocrystalline diamond coatings deposited by HFCVD on cemented carbide grades with different cobalt contents. The test duration was limited to a maximum of 300 s.</figcaption></figure>

Critical strain of CVD-diamond-coatings. In order to determine the limiting deformation of the CVD coatings, hot compression tests were carried out using the WC-Co alloys EMT100 and EMT609 as substrate materials. The applied strain was derived from the previous hot compression tests of uncoated samples. Here, an average fracture strain of 0.0333 and 0.0325 was measured for EMT100 and EMT609, respectively. Therefore, the targeted strain values for the three iterations were set to 0.017 (50 % fracture strain), 0.025 (75 % fracture strain) and 0.029 (87,5 % fracture strain) for the three iteration stages. The first substrate material to be tested was EMT100. In the first iteration stage, no defects of the CVD coating were observed. Regardless of the coating configuration (morphology/thickness) or substrate pre-treatment, no visually recognizable damage to the coating occurred. Figure 4a illustrates a specimen, which was tested in the first iteration stage. In the second iteration stage, repeated failure of the substrate material occurred even though the applied strain was below the determined fracture strain. An example is shown in Figure 4b, where the specimen is split in half. Due to the early fracture of the WC-Co substrate, the
designated strain could not be applied in all trials of the second iteration stage. However, no damage to the coatings caused by the compression was detected either. The microscopic examination of the fracture surface did reveal no damage to the interface between diamond coating and substrate (see Figure 5a). An exception is the microcrystalline, 12 µm thick diamond coating on ground substrate material. Here, a large area of spalling of the coating could be seen after the specimen failed (Figure 4c). In Figure 5b the delamination of the coating is visible. The gap between substrate (left) and diamond film (center) clearly indicates the separation of the interface. Furthermore, at the top of the picture a piece of the coating is missing. In the third iteration stage, the desired strain was achieved in only a few trials. In most cases, the WC-Co failed at strains close to 0.025. The frequent fracture in form of longitudinal splitting through the center of the specimens suggests the laser engraving on the cylinder front face acts as a crack initiation point. Unstable crack growth is known to be preceded by crack initiation and to be characteristic for brittle materials under mechanical load [3]. Examination of the broken specimens showed that the fracture always extends through digits of the laser engraving. Multiple additional tests with reference specimens (same batch of substrate material, different coating systems, without engraving) confirmed laser engraving as the cause of premature failure of the specimens. In the tests in which the specified strain of 0.029 was achieved, layer failure occurred in the form of large-area ablation (see Figure 4d). This suggests the critical strain for coating failure is in the range 0.025-0.029, which is above the compressive yield strain of the substrate material. This is not critical for the intended application of CVD diamond coating on dies for aluminum or magnesium extrusion, as plastic deformation of the die must be avoided.

<figure><img src="image_644.png" /><figcaption>Figure 4: Observed failure modes: a) no failure, b) ablation without substrate failure, c) substrate failure without ablation and d) ablation with substrate failure.</figcaption></figure>

<figure><img src="image_645.png" /><figcaption>Figure 5: Interface of substrate and diamond coating after specimen fracture with microcrystalline coating on ground substrate. a) fully intact interface for the 6 µm thick coating, b) delamination of the 12 µm thick coating.</figcaption></figure>
designated strain could not be applied in all trials of the second iteration stage. However, no damage to the coatings caused by the compression was detected either. The microscopic examination of the fracture surface did reveal no damage to the interface between diamond coating and substrate (see Figure 5a). An exception is the microcrystalline, 12 µm thick diamond coating on ground substrate material. Here, a large area of spalling of the coating could be seen after the specimen failed (Figure 4c). In Figure 5b the delamination of the coating is visible. The gap between substrate (left) and diamond film (center) clearly indicates the separation of the interface. Furthermore, at the top of the picture a piece of the coating is missing. In the third iteration stage, the desired strain was achieved in only a few trials. In most cases, the WC-Co failed at strains close to 0.025. The frequent fracture in form of longitudinal splitting through the center of the specimens suggests the laser engraving on the cylinder front face acts as a crack initiation point. Unstable crack growth is known to be preceded by crack initiation and to be characteristic for brittle materials under mechanical load [3]. Examination of the broken specimens showed that the fracture always extends through digits of the laser engraving. Multiple additional tests with reference specimens (same batch of substrate material, different coating systems, without engraving) confirmed laser engraving as the cause of premature failure of the specimens. In the tests in which the specified strain of 0.029 was achieved, layer failure occurred in the form of large-area ablation (see Figure 4d). This suggests the critical strain for coating failure is in the range 0.025-0.029, which is above the compressive yield strain of the substrate material. This is not critical for the intended application of CVD diamond coating on dies for aluminum or magnesium extrusion, as plastic deformation of the die must be avoided.

<figure><img src="image_655.png" /><figcaption>Figure 4: Observed failure modes: a) no failure, b) ablation without substrate failure, c) substrate failure without ablation and d) ablation with substrate failure.</figcaption></figure>

<figure><img src="image_656.png" /><figcaption>Figure 5: Interface of substrate and diamond coating after specimen fracture with microcrystalline coating on ground substrate. a) fully intact interface for the 6 µm thick coating, b) delamination of the 12 µm thick coating.</figcaption></figure>
The second substrate material on which the CVD-diamond coating was applied is EMT609. With this WC-Co alloy, fractures of specimens were already observed during the first iteration stage. As described above, this is most likely due to the laser engraving and its action as a crack initiation point. The reason why material failure occurs at lower strains than with EMT100 is due to the increased strength of EMT609 [13]. With the same strain, higher stresses are present in the higher-strength material. Conversely, the strain of the higher strength-material is lower for the same load. Accordingly, EMT609 in combination with the weakening caused by the laser engraving fails at lower strains compared EMT100. Regardless of this, no damage to the CVD coating was visible after the tests on the first iteration stage. In the second iteration stage, frequent specimen failure was observed and only very occasionally the targeted 75% of fracture strain was achieved. In most cases, the samples failed at strains of 0.018-0.020. Nevertheless, the CVD coatings remained intact. Only the microcrystalline, 12µm thick diamond coating on ground substrate showed distinct layer failure in form of large area spalling. Due to the increased layer thickness, the coating exhibits reduced flexibility and can therefore tolerate less deformation. The tests of the third iteration were carried out as planned, but as in the second iteration stage no strains above 0.020 were achieved.

## Summary

In this study, the suitability of diamond films deposited by HFCVD on different types of WC-Co substrates was investigated regarding its application on extrusion dies. The mechanical material properties of four WC-Co alloys (6-12 % cobalt, fine and ultrafine grain) were determined at elevated temperatures (300-500 °C) and two WC-Co materials were selected for further investigations. The alloys EMT100 and EMT609 were used to investigate the adhesion and critical strain of the diamond coating. The knowledge gained can be summarized as follows:

- Compressive yield strain and compressive yield strength at elevated temperatures of the investigated WC-Co alloys are directly proportional to their room temperature hardness.

- The combination of moderate cobalt content and ultrafine grain of the EMT609 alloy offers a good balance of toughness and strength.

- For coatings on EMT100, a compressive strain of 0.025-0.029 was identified as critical for coating failure.

- Due to premature specimen failure, coatings on EMT609 could only be subjected to a strain of up to 0.020. The coatings remained undamaged proving excellent adhesion levels of the deposited diamond films.

- The nanocrystalline diamond films show excellent stability in hot compression tests for both film thicknesses (6 and 12 µm) independent on the type of surface finishing (ground or roughened by sandblasting). The parameter combination of microcrystalline, 12 µm thick coating on ground substrate seems to be unfavorable for the critical strain of the diamond coating, as cumulative ablation was observed.

## References

[1] M. Bauser, G. Sauer, K. Siegert, Strangpressen: Aluminium Fachbuchreihe, Aluminium-Verlag, Düsseldorf, 2001. ISBN 3410221816

[2] H.J. Weintraud, Neue Werkstoffe für Strangpressmatrizen, in: J. Baumgarten (Ed.), Strangpressen, DGM-Informationsgesellschaft m.b.H., Oberursel, 1989. ISBN 3-88355-160-0

[3] Upadhyaya, G.S., Cemented Tungsten Carbides – Production, Properties, and Testing, William Andrew, Norwich, 1998. ISBN 978-0-8155-1417-6. https://doi.org/10.1016/B978-081551417-6.50002-2

[4] A.G. Guy, G. Petzow, Metallkunde für Ingenieure, in R. Sewig, D. Roess (Eds.), Technisch-Physikalische Sammlung Band 7, Akademische Verlagsgesellschaft, Frankfurt am Main, 1970.
The second substrate material on which the CVD-diamond coating was applied is EMT609. With this WC-Co alloy, fractures of specimens were already observed during the first iteration stage. As described above, this is most likely due to the laser engraving and its action as a crack initiation point. The reason why material failure occurs at lower strains than with EMT100 is due to the increased strength of EMT609 [13]. With the same strain, higher stresses are present in the higher-strength material. Conversely, the strain of the higher strength-material is lower for the same load. Accordingly, EMT609 in combination with the weakening caused by the laser engraving fails at lower strains compared EMT100. Regardless of this, no damage to the CVD coating was visible after the tests on the first iteration stage. In the second iteration stage, frequent specimen failure was observed and only very occasionally the targeted 75% of fracture strain was achieved. In most cases, the samples failed at strains of 0.018-0.020. Nevertheless, the CVD coatings remained intact. Only the microcrystalline, 12µm thick diamond coating on ground substrate showed distinct layer failure in form of large area spalling. Due to the increased layer thickness, the coating exhibits reduced flexibility and can therefore tolerate less deformation. The tests of the third iteration were carried out as planned, but as in the second iteration stage no strains above 0.020 were achieved.

## Summary

In this study, the suitability of diamond films deposited by HFCVD on different types of WC-Co substrates was investigated regarding its application on extrusion dies. The mechanical material properties of four WC-Co alloys (6-12 % cobalt, fine and ultrafine grain) were determined at elevated temperatures (300-500 °C) and two WC-Co materials were selected for further investigations. The alloys EMT100 and EMT609 were used to investigate the adhesion and critical strain of the diamond coating. The knowledge gained can be summarized as follows:

- Compressive yield strain and compressive yield strength at elevated temperatures of the investigated WC-Co alloys are directly proportional to their room temperature hardness.

- The combination of moderate cobalt content and ultrafine grain of the EMT609 alloy offers a good balance of toughness and strength.

- For coatings on EMT100, a compressive strain of 0.025-0.029 was identified as critical for coating failure.

- Due to premature specimen failure, coatings on EMT609 could only be subjected to a strain of up to 0.020. The coatings remained undamaged proving excellent adhesion levels of the deposited diamond films.

- The nanocrystalline diamond films show excellent stability in hot compression tests for both film thicknesses (6 and 12 µm) independent on the type of surface finishing (ground or roughened by sandblasting). The parameter combination of microcrystalline, 12 µm thick coating on ground substrate seems to be unfavorable for the critical strain of the diamond coating, as cumulative ablation was observed.

## References

[1] M. Bauser, G. Sauer, K. Siegert, Strangpressen: Aluminium Fachbuchreihe, Aluminium-Verlag, Düsseldorf, 2001. ISBN 3410221816

[2] H.J. Weintraud, Neue Werkstoffe für Strangpressmatrizen, in: J. Baumgarten (Ed.), Strangpressen, DGM-Informationsgesellschaft m.b.H., Oberursel, 1989. ISBN 3-88355-160-0

[3] Upadhyaya, G.S., Cemented Tungsten Carbides – Production, Properties, and Testing, William Andrew, Norwich, 1998. ISBN 978-0-8155-1417-6. https://doi.org/10.1016/B978-081551417-6.50002-2

[4] A.G. Guy, G. Petzow, Metallkunde für Ingenieure, in R. Sewig, D. Roess (Eds.), Technisch-Physikalische Sammlung Band 7, Akademische Verlagsgesellschaft, Frankfurt am Main, 1970.
[5] P. Kazanowski, A Study of Multi-material Die Insert Failure Mechanism in Aluminum Extrusion, Conf. 10th International Aluminium Extrusion Technology Seminar ET (2012), Miami.

[6] M.P. Clode, T. Sheppard, Formation of die lines during extrusion of AA6063, Mater. Sci. Technol., 6(8) (1990), 755-763. https://doi.org/10.1179/mst.1990.6.8.755

[7] K. Müller, Grundlagen des Strangpressens: Verfahren, Anlagen, Werkstoffe, Werkzeuge, expert verlag, Renningen, 2003. ISBN 3-81692-018-7

[8] A.K. Chattopadhyay, P. Roy, A. Gosch, S.K. Sarangi, Wettability and machinability study of pure aluminium towards uncoated and coated carbide cutting tool inserts, Surf. Coat. Technol. 203 (2009), 954-951. https://doi.org/10.1016/j.surfcoat.2008.08.047

[9] S. A. Linnik, A. V. Gaidaichuk, and V. V. Okhotnikov, Influence of Cobalt on the Adhesion Strength of Polycrystalline Diamond Coatings on WC-Co Hard Alloys, Technical Physics, 63(2) (2018), 206-210, https://doi.org/10.1134/S1063784218020226

[10] H.E. Exner, J. Gurland, A review of parameters influencing some mechanical properties of tungsten carbide-cobalt alloys, Powder Metall., 13:25 (1970), 13-31. https://doi.org/10.1179/pom.1970.13.25.002

[11] A.V. Shatov, S.S. Ponomarev, S.A. Firstov, Fracture and Strength of Hard metals at Room Temperature, in: V.K. Sarin (Ed.), Comprehensive Hard Metals, Elsevier, 2014, pp. 301-343. https://doi.org/10.1016/B978-0-08-096527-7.00010-6

[12] P. Ettmayer, Hardmetals and Cermets, Annu. Rev. Mater. Sci. 19:1 (1989), 145-164. https://doi.org/10.1146/annurev.ms.19.080189.001045

[13] EXTRAMET AG, Delivery Program 09/2022, Plaffeien, 2022. https://www.extramet.ch/wp-content/uploads/2022/08/EXTRAMET_Lieferprogramm-2022 www_202208.pdf
[5] P. Kazanowski, A Study of Multi-material Die Insert Failure Mechanism in Aluminum Extrusion, Conf. 10th International Aluminium Extrusion Technology Seminar ET (2012), Miami.

[6] M.P. Clode, T. Sheppard, Formation of die lines during extrusion of AA6063, Mater. Sci. Technol., 6(8) (1990), 755-763. https://doi.org/10.1179/mst.1990.6.8.755

[7] K. Müller, Grundlagen des Strangpressens: Verfahren, Anlagen, Werkstoffe, Werkzeuge, expert verlag, Renningen, 2003. ISBN 3-81692-018-7

[8] A.K. Chattopadhyay, P. Roy, A. Gosch, S.K. Sarangi, Wettability and machinability study of pure aluminium towards uncoated and coated carbide cutting tool inserts, Surf. Coat. Technol. 203 (2009), 954-951. https://doi.org/10.1016/j.surfcoat.2008.08.047

[9] S. A. Linnik, A. V. Gaidaichuk, and V. V. Okhotnikov, Influence of Cobalt on the Adhesion Strength of Polycrystalline Diamond Coatings on WC-Co Hard Alloys, Technical Physics, 63(2) (2018), 206-210, https://doi.org/10.1134/S1063784218020226

[10] H.E. Exner, J. Gurland, A review of parameters influencing some mechanical properties of tungsten carbide-cobalt alloys, Powder Metall., 13:25 (1970), 13-31. https://doi.org/10.1179/pom.1970.13.25.002

[11] A.V. Shatov, S.S. Ponomarev, S.A. Firstov, Fracture and Strength of Hard metals at Room Temperature, in: V.K. Sarin (Ed.), Comprehensive Hard Metals, Elsevier, 2014, pp. 301-343. https://doi.org/10.1016/B978-0-08-096527-7.00010-6

[12] P. Ettmayer, Hardmetals and Cermets, Annu. Rev. Mater. Sci. 19:1 (1989), 145-164. https://doi.org/10.1146/annurev.ms.19.080189.001045

[13] EXTRAMET AG, Delivery Program 09/2022, Plaffeien, 2022. https://www.extramet.ch/wp-content/uploads/2022/08/EXTRAMET_Lieferprogramm-2022 www_202208.pdf
# Prediction of the microstructure evolution during the friction stir extrusion of a AA6061 aluminum alloy

BOCCHI Sara<fnref n="1" />,*, NEGOZIO Marco<fnref n="2" />,ᵇ, GIARDINI Claudio<fnref n="1" />,ᶜ, DONATI Lorenzo<fnref n="3" />,ᵈ

**Keywords:** Friction Stir Extrusion, FEM, Recrystallization, Process Sustainability, AA6XXX

**Abstract.** In recent years, the development of Friction Stir Extrusion (FSE) simulation models becomes crucial for gaining a deeper understanding of its underlying physics. Concurrently, there is a demand for precise control over the microstructure evolution of aluminum alloy extruded profiles, given its substantial impact on mechanical properties. Despite this, the relationship between process parameters and the evolution of grain structure remains insufficiently understood. In this context, a Lagrangian approach was established to simulate the FSE process, utilizing the commercial software DEFORM™ 3D. This research involved the investigation of the impact of various process parameters, such as rotational and descent tool speeds, on the occurrence of bonding phenomena, while considering both thermal and stress conditions. Furthermore, an innovative model originally developed for traditionally extruded components was implemented in a customized Fortran post-processing routine to investigate and predict the recrystallization behavior in the FSE of AA6061 aluminum alloy.

## Introduction

In recent years, the European Commission has been actively pursuing sustainable growth strategies with a focus on competitive energy resources, taking into account both environmental impact and economic considerations [1]. A significant aspect of this endeavor involves addressing the challenges associated with the extensive waste generated by mechanical processes in industries such as metalworking. Conventional recycling methods often lead to increased energy consumption and difficulties in maintaining the desired alloy composition [2,3].

To overcome these challenges, the Welding Institute introduced Friction Stir Extrusion in 1993 as a novel metal shavings recycling process. FSE, characterized by solid-state recycling and high energy efficiency, has emerged as a promising technology, utilizing friction-generated heat for plastic deformation and extrusion of metal chips [4]. This process has demonstrated a remarkable reduction in energy consumption, requiring less than 15% of the energy used in conventional aluminum alloy melting processes [5]. In parallel with these advancements, it has become apparent that a crucial factor in optimization is understanding the intricate relationship between process parameters and the resulting physical outcomes [6]. Furthermore, the extrusion industry has long acknowledged the substantial attention given to the microstructure of aluminum alloy extruded profiles. The complex connection between process parameters and the evolution of grain structure presents challenges for extrusion companies striving to enhance product performance. So far, the scientific community has been actively engaged in investigating the recrystallization of aluminum

<figure><img src="image_646.png" /></figure>

<footnote n="1">¹ Department of Management, Information and Production Engineering, University of Bergamo,
Via Pasubio 7/b, Dalmine (BG), Italy</footnote>

<footnote n="2">² University of Parma - Department of Engineering and Architecture, Parco Area delle Scienze,
181/A, 43124 Parma, Italy</footnote>

<footnote n="3">³ DIN Department of Industrial Engineering – University of Bologna,
Viale Risorgimento 2, 40136, Bologna, Italy</footnote>

<footnote>ªsara.bocchi@unibg.it, ᵇmarco.negozio@unipr.it, ᶜclaudio.giardini@unibg.it, ᵈl.donati@unibo.it</footnote>
# Prediction of the microstructure evolution during the friction stir extrusion of a AA6061 aluminum alloy

BOCCHI Sara<fnref n="1" />,*, NEGOZIO Marco<fnref n="2" />,ᵇ, GIARDINI Claudio<fnref n="1" />,ᶜ, DONATI Lorenzo<fnref n="3" />,ᵈ

**Keywords:** Friction Stir Extrusion, FEM, Recrystallization, Process Sustainability, AA6XXX

**Abstract.** In recent years, the development of Friction Stir Extrusion (FSE) simulation models becomes crucial for gaining a deeper understanding of its underlying physics. Concurrently, there is a demand for precise control over the microstructure evolution of aluminum alloy extruded profiles, given its substantial impact on mechanical properties. Despite this, the relationship between process parameters and the evolution of grain structure remains insufficiently understood. In this context, a Lagrangian approach was established to simulate the FSE process, utilizing the commercial software DEFORM™ 3D. This research involved the investigation of the impact of various process parameters, such as rotational and descent tool speeds, on the occurrence of bonding phenomena, while considering both thermal and stress conditions. Furthermore, an innovative model originally developed for traditionally extruded components was implemented in a customized Fortran post-processing routine to investigate and predict the recrystallization behavior in the FSE of AA6061 aluminum alloy.

## Introduction

In recent years, the European Commission has been actively pursuing sustainable growth strategies with a focus on competitive energy resources, taking into account both environmental impact and economic considerations [1]. A significant aspect of this endeavor involves addressing the challenges associated with the extensive waste generated by mechanical processes in industries such as metalworking. Conventional recycling methods often lead to increased energy consumption and difficulties in maintaining the desired alloy composition [2,3].

To overcome these challenges, the Welding Institute introduced Friction Stir Extrusion in 1993 as a novel metal shavings recycling process. FSE, characterized by solid-state recycling and high energy efficiency, has emerged as a promising technology, utilizing friction-generated heat for plastic deformation and extrusion of metal chips [4]. This process has demonstrated a remarkable reduction in energy consumption, requiring less than 15% of the energy used in conventional aluminum alloy melting processes [5]. In parallel with these advancements, it has become apparent that a crucial factor in optimization is understanding the intricate relationship between process parameters and the resulting physical outcomes [6]. Furthermore, the extrusion industry has long acknowledged the substantial attention given to the microstructure of aluminum alloy extruded profiles. The complex connection between process parameters and the evolution of grain structure presents challenges for extrusion companies striving to enhance product performance. So far, the scientific community has been actively engaged in investigating the recrystallization of aluminum

<figure><img src="image_657.png" /></figure>

<footnote n="1">¹ Department of Management, Information and Production Engineering, University of Bergamo,
Via Pasubio 7/b, Dalmine (BG), Italy</footnote>

<footnote n="2">² University of Parma - Department of Engineering and Architecture, Parco Area delle Scienze,
181/A, 43124 Parma, Italy</footnote>

<footnote n="3">³ DIN Department of Industrial Engineering – University of Bologna,
Viale Risorgimento 2, 40136, Bologna, Italy</footnote>

<footnote>ªsara.bocchi@unibg.it, ᵇmarco.negozio@unipr.it, ᶜclaudio.giardini@unibg.it, ᵈl.donati@unibo.it</footnote>
alloys during the extrusion process. Efforts have been directed towards developing reliable models for recrystallization prediction, essential for controlling the final microstructure at the die design stage without costly experimental trial and error [6,7].

In the traditionally extruded pieces, the recrystallized structure, characterized by equiaxed grains, offers an avenue for optimizing product performance if appropriately controlled. Two main recrystallization mechanisms, dynamic recrystallization (DRX) and static recrystallization (SRX), have been identified, each influencing the microstructure during and after the extrusion process [8,9]. Despite numerous studies on recrystallization kinetics, limitations persist in the validation of models against industrial-scale extrusions with complex geometry profiles.

In response to this gap, this paper investigates the FSE process considering an already developed and optimized Finite Element simulation performed within the DEFORM™ 3D code. This model explores the influence of rotational tool speed and descent speed on process conditions, such as attained temperatures and the final volume of the workpiece. Additionally, stress conditions resulting from the simulations were analyzed using the Piwnik and Plata criterion to investigate the bonding phenomenon [10]. The study also aims to provide a comprehensive understanding of the recrystallization mechanisms involved. A recrystallization model, originally developed and validated for the traditional extrusion process [11], has been adapted for the FSE process. The ultimate objective is to propose a reliable model capable of accurately predicting the microstructural behavior in the AA6061 aluminum alloy Friction Stir Extruded (FSEed) profiles.

## Procedure

A 3D Finite Element Lagrangian model was created using DEFORM™ 3D software to simulate the extrusion process. The tool and container were modeled as rigid bodies using AISI-1043 steel, while the aluminum scraps were treated as a porous workpiece of Aluminum-6061. The initial compaction phase was excluded for computational efficiency, and the initial density was set to 78% of the base aluminum, based on experimental tests, as also reported in [12–14]. Material flow stress data were sourced from the DEFORM™ 3D library for aluminum and steel, following both the law reported in Eq. 1.

$$ \bar{\sigma} = \bar{\sigma}(\bar{\varepsilon}, \dot{\bar{\varepsilon}}, T) \qquad (1) $$

The steel tool and container were considered rigid due to the significant difference in yield strength compared to aluminum. The porous workpiece was meshed with 50,000 tetrahedral elements (Fig. 1).

<figure><img src="image_647.png" /><figcaption>Figure 1: Geometry of the all simulated structure.</figcaption></figure>

Thermal parameters were kept constant (Table 1), and simulations were conducted by varying tool rotational speed and descent speed to understand the bonding phenomenon (Table 2). The stop criterion was a tool displacement of 5 mm in the –Z direction.
alloys during the extrusion process. Efforts have been directed towards developing reliable models for recrystallization prediction, essential for controlling the final microstructure at the die design stage without costly experimental trial and error [6,7].

In the traditionally extruded pieces, the recrystallized structure, characterized by equiaxed grains, offers an avenue for optimizing product performance if appropriately controlled. Two main recrystallization mechanisms, dynamic recrystallization (DRX) and static recrystallization (SRX), have been identified, each influencing the microstructure during and after the extrusion process [8,9]. Despite numerous studies on recrystallization kinetics, limitations persist in the validation of models against industrial-scale extrusions with complex geometry profiles.

In response to this gap, this paper investigates the FSE process considering an already developed and optimized Finite Element simulation performed within the DEFORM™ 3D code. This model explores the influence of rotational tool speed and descent speed on process conditions, such as attained temperatures and the final volume of the workpiece. Additionally, stress conditions resulting from the simulations were analyzed using the Piwnik and Plata criterion to investigate the bonding phenomenon [10]. The study also aims to provide a comprehensive understanding of the recrystallization mechanisms involved. A recrystallization model, originally developed and validated for the traditional extrusion process [11], has been adapted for the FSE process. The ultimate objective is to propose a reliable model capable of accurately predicting the microstructural behavior in the AA6061 aluminum alloy Friction Stir Extruded (FSEed) profiles.

## Procedure

A 3D Finite Element Lagrangian model was created using DEFORM™ 3D software to simulate the extrusion process. The tool and container were modeled as rigid bodies using AISI-1043 steel, while the aluminum scraps were treated as a porous workpiece of Aluminum-6061. The initial compaction phase was excluded for computational efficiency, and the initial density was set to 78% of the base aluminum, based on experimental tests, as also reported in [12–14]. Material flow stress data were sourced from the DEFORM™ 3D library for aluminum and steel, following both the law reported in Eq. 1.

$$ \bar{\sigma} = \bar{\sigma}(\bar{\varepsilon}, \dot{\bar{\varepsilon}}, T) \qquad (1) $$

The steel tool and container were considered rigid due to the significant difference in yield strength compared to aluminum. The porous workpiece was meshed with 50,000 tetrahedral elements (Fig. 1).

<figure><img src="image_658.png" /><figcaption>Figure 1: Geometry of the all simulated structure.</figcaption></figure>

Thermal parameters were kept constant (Table 1), and simulations were conducted by varying tool rotational speed and descent speed to understand the bonding phenomenon (Table 2). The stop criterion was a tool displacement of 5 mm in the –Z direction.
<table><caption>Table 1: Thermal parameters used in the simulations [12].</caption><thead><tr><td>Parameters</td><td>Value</td></tr></thead><tbody><tr><td>Friction coefficient aluminum-tool</td><td>0.60</td></tr><tr><td>Thermal conductivity [N/(mm·s·°C)]</td><td>450</td></tr><tr><td>Aluminum emissivity</td><td>0.25</td></tr><tr><td>Heat transfer coefficient aluminum-tool [N/(mm·s·°C)]</td><td>11.00</td></tr><tr><td>Heat exchange with the environment [N/(mm·s·°C)]</td><td>0.02</td></tr><tr><td>Mechanical conversion to heat</td><td>0.80</td></tr></tbody></table>

<table><caption>Table 2: Combinations between the rotational speed (S) and the descent speed (F) considered for the simulations [12].</caption><thead><tr><td>Combination</td><td>Rotational speed (S) [rpm]</td><td>Descent speed (F) [mm/s]</td></tr></thead><tbody><tr><td>#1</td><td>400</td><td>1</td></tr><tr><td>#2</td><td>400</td><td>3</td></tr><tr><td>#3</td><td>800</td><td>1</td></tr><tr><td>#4</td><td>800</td><td>3</td></tr></tbody></table>

According to Piwnik and Plata criterion, material bonding is established when the parameter w, characterized as the integrated ratio over time between the pressure and the effective stress exerted on the material, attains a predetermined threshold value known as w<sub>lim</sub>. This threshold is determined as a function of the temperature. To compute w and w<sub>lim</sub> for every node at each step, a dedicated Fortran routine was specifically created and integrated into the simulation model. This ensured the automated calculation of these parameters and their relationship within the model. Moreover, a novel model initially designed for conventionally extruded components was incorporated into a tailored Fortran post-processing routine. This adaptation aimed to explore and anticipate the recrystallization behavior during the FSE of AA6061 aluminum alloy.

## Modeling

The recrystallization model was taken from the work of Negozio et al. [15]. In this study, the model was developed and implemented in a subroutine for the assessment of dynamic and static recrystallization in post-processing to the Finite Element Method simulation of the extrusion process. To assess the microstructural evolution, the calculation of the Zener-Hollomon Z and sub-grain size δ parameters was carried out according to [16].

$$Z = \dot{\epsilon} \exp \left( \frac{Q}{RT} \right) \qquad (2)$$

$$\frac{1}{\delta} = C (\ln Z)^n \qquad (3)$$

where $\dot{\epsilon}$ is the maximum strain rate reached by each point during the investigated plastic deformation process, T is the temperature, n=5.577, C=3.36x10<sup>-9</sup> m<sup>-1</sup>, Q is the activation energy of the AA6061 aluminum alloy (196000 J/mol*K [17]), R is the universal gas constant (8.341 J/mol).

The energy responsible for activating the kinetics of static recrystallization is referred to as Stored Energy Pd and is calculated in accordance with [6]:

$$Pd = \frac{Gb^2}{10} \left[ \rho_i(1-\ln(10b\rho_i^{0.5})) + \frac{2\theta}{b\delta} * \left(1 + \ln\left(\frac{\theta c}{\theta}\right)\right) \right] \qquad (4)$$
<table><caption>Table 1: Thermal parameters used in the simulations [12].</caption><thead><tr><td>Parameters</td><td>Value</td></tr></thead><tbody><tr><td>Friction coefficient aluminum-tool</td><td>0.60</td></tr><tr><td>Thermal conductivity [N/(mm·s·°C)]</td><td>450</td></tr><tr><td>Aluminum emissivity</td><td>0.25</td></tr><tr><td>Heat transfer coefficient aluminum-tool [N/(mm·s·°C)]</td><td>11.00</td></tr><tr><td>Heat exchange with the environment [N/(mm·s·°C)]</td><td>0.02</td></tr><tr><td>Mechanical conversion to heat</td><td>0.80</td></tr></tbody></table>

<table><caption>Table 2: Combinations between the rotational speed (S) and the descent speed (F) considered for the simulations [12].</caption><thead><tr><td>Combination</td><td>Rotational speed (S) [rpm]</td><td>Descent speed (F) [mm/s]</td></tr></thead><tbody><tr><td>#1</td><td>400</td><td>1</td></tr><tr><td>#2</td><td>400</td><td>3</td></tr><tr><td>#3</td><td>800</td><td>1</td></tr><tr><td>#4</td><td>800</td><td>3</td></tr></tbody></table>

According to Piwnik and Plata criterion, material bonding is established when the parameter w, characterized as the integrated ratio over time between the pressure and the effective stress exerted on the material, attains a predetermined threshold value known as w<sub>lim</sub>. This threshold is determined as a function of the temperature. To compute w and w<sub>lim</sub> for every node at each step, a dedicated Fortran routine was specifically created and integrated into the simulation model. This ensured the automated calculation of these parameters and their relationship within the model. Moreover, a novel model initially designed for conventionally extruded components was incorporated into a tailored Fortran post-processing routine. This adaptation aimed to explore and anticipate the recrystallization behavior during the FSE of AA6061 aluminum alloy.

## Modeling

The recrystallization model was taken from the work of Negozio et al. [15]. In this study, the model was developed and implemented in a subroutine for the assessment of dynamic and static recrystallization in post-processing to the Finite Element Method simulation of the extrusion process. To assess the microstructural evolution, the calculation of the Zener-Hollomon Z and sub-grain size δ parameters was carried out according to [16].

$$Z = \dot{\epsilon} \exp \left( \frac{Q}{RT} \right) \qquad (2)$$

$$\frac{1}{\delta} = C (\ln Z)^n \qquad (3)$$

where $\dot{\epsilon}$ is the maximum strain rate reached by each point during the investigated plastic deformation process, T is the temperature, n=5.577, C=3.36x10<sup>-9</sup> m<sup>-1</sup>, Q is the activation energy of the AA6061 aluminum alloy (196000 J/mol*K [17]), R is the universal gas constant (8.341 J/mol).

The energy responsible for activating the kinetics of static recrystallization is referred to as Stored Energy Pd and is calculated in accordance with [6]:

$$P_d = \frac{Gb^2}{10} \left[ \rho_i (1 - \ln(10b\rho_i^{0.5})) + \frac{2\theta}{b\delta} * \left( 1 + \ln\left(\frac{\theta c}{\theta}\right) \right) \right] \qquad (4)$$
where b the Burgers vector (2.86x10⁻¹⁰ m [18]), G is the material shear modulus (2.05x10¹⁰ Pa [18]), ρᵢ the dislocation density, δ the sub-grain size, θ the misorientation angle and θc the misorientation angle limit (15°) [15].

The experimental analysis revealed a completely fibrous microstructure, without showing evidence of static recrystallization. Dynamic recrystallization modeling was conducted following the guidelines provided in [15]. The evolution of the grain size subjected to deformation was described based on the definition of two quantities, dₜ (grain thickness) and dₜ (grain length). These quantities were calculated as follows [15]:

$$d_t = (d_0 - 2.5\delta_{ss})(k_1)^{\bar{\epsilon}} + 2.5\delta_{ss} \quad (5)$$

$$d_l = k_2 \bar{\epsilon}^2 - k_3 \bar{\epsilon} + d_0 \quad \text{if } \bar{\epsilon} < \epsilon_p \quad (6)$$

$$d_l = k_4 \bar{\epsilon}^{-m} + 10\delta_{ss} \quad \text{if } \bar{\epsilon} > \epsilon_p \quad (7)$$

where εₚ is the critical strain level for the pinch-off occurrence [17], which is considered as value of 3 [19], δₐ₅ is the subgrain size at the steady-state condition (δₐ₅ = 8.4 µm) [19]) and m, k₁, k₂, k₃, k₄ are material constants identified in [19] as m = 4.75, k₁ = 0.4, k₂ = 85.192, k₃ = 14.88, k₄ = 1.68 * 10⁵.

## Results and discussion

When the occurrence of bonding according to the Piwnik and Plata criterion is verified, the Fortran variable named "Welding" shifts from 0 to 1. In such cases, the element undergoing this condition is visually highlighted in red. Conversely, blue indicates regions where stress conditions and temperature distribution are insufficient for the bonding of scraps. This approach facilitates pinpointing the occurrence and location of material bonding, as depicted in Fig. 2.

<figure><img src="image_648.png" /><figcaption>Figure 2: Red and blue areas in the FSEed object obtained with S=800 rpm, F=1 mm/s and Z=2.5 mm.</figcaption></figure>

The examination of the simulation results revealed three global conditions: bonding, non-bonding, and limit conditions, across various combinations of operational parameters. Differentiating between instances where metal scraps have achieved complete bonding and cases where bonding was not accomplished before extrusion is crucial. This distinction is essential not only for identifying the presence of defects or voids within the extruded object but also for ensuring a comprehensive understanding of the bonding process. The limit condition involves the simultaneous presence of fully bonded material and still-porous metal scraps: this condition allows
where b the Burgers vector (2.86x10⁻¹⁰ m [18]), G is the material shear modulus (2.05x10¹⁰ Pa [18]), ρᵢ the dislocation density, δ the sub-grain size, θ the misorientation angle and θc the misorientation angle limit (15°) [15].

The experimental analysis revealed a completely fibrous microstructure, without showing evidence of static recrystallization. Dynamic recrystallization modeling was conducted following the guidelines provided in [15]. The evolution of the grain size subjected to deformation was described based on the definition of two quantities, dᵡ (grain thickness) and dᵢ (grain length). These quantities were calculated as follows [15]:

$$d_t = (d_0 - 2.5\delta_{ss})(k_1)^{\bar{\epsilon}} + 2.5\delta_{ss} \quad (5)$$

$$d_l = k_2 \bar{\epsilon}^2 - k_3 \bar{\epsilon} + d_0 \quad \text{if } \bar{\epsilon} < \epsilon_p \quad (6)$$

$$d_l = k_4 \bar{\epsilon}^{-m} + 10\delta_{ss} \quad \text{if } \bar{\epsilon} > \epsilon_p \quad (7)$$

where εₚ is the critical strain level for the pinch-off occurrence [17], which is considered as value of 3 [19], δₐ₅ is the subgrain size at the steady-state condition (δₐ₅ = 8.4 µm) [19]) and m, k₁, k₂, k₃, k₄ are material constants identified in [19] as m = 4.75, k₁ = 0.4, k₂ = 85.192, k₃ = 14.88, k₄ = 1.68 * 10⁵.

## Results and discussion

When the occurrence of bonding according to the Piwnik and Plata criterion is verified, the Fortran variable named "Welding" shifts from 0 to 1. In such cases, the element undergoing this condition is visually highlighted in red. Conversely, blue indicates regions where stress conditions and temperature distribution are insufficient for the bonding of scraps. This approach facilitates pinpointing the occurrence and location of material bonding, as depicted in Fig. 2.

<figure><img src="image_659.png" /><figcaption>Figure 2: Red and blue areas in the FSEed object obtained with S=800 rpm, F=1 mm/s and Z=2.5 mm.</figcaption></figure>

The examination of the simulation results revealed three global conditions: bonding, non-bonding, and limit conditions, across various combinations of operational parameters. Differentiating between instances where metal scraps have achieved complete bonding and cases where bonding was not accomplished before extrusion is crucial. This distinction is essential not only for identifying the presence of defects or voids within the extruded object but also for ensuring a comprehensive understanding of the bonding process. The limit condition involves the simultaneous presence of fully bonded material and still-porous metal scraps: this condition allows
for the production of extruded pieces with a solid external surface but internal sections of unbounded material. Identifying this scenario is essential, as it could lead to significant issues, resulting in a more fragile behavior than a completely solid FSEed sample.

Addressing these challenges requires precise definition of the boundary condition zone and the selection of parameter combinations guaranteed to fall within this range. This determination is facilitated by consulting the diagram presented in Fig. 3, which illustrates the summarized experimental results in terms of bonded and non-bonded specimens. The red cross designates the non-bonding zone, the green dot signifies the bonding zone, and the yellow triangle represents the limit condition zone.

<figure><img src="image_649.png" /><figcaption>Figure 3: Bonding conditions as a function of the process parameters.</figcaption></figure>

From the simulations point of view, the bonding, no bonding and limit conditions are reported in Fig. 4. The scenario where metal scraps are fully bonded (Fig. 4a) contrasts notably with instances where bonding fails to occur before extrusion (Fig. 4c).

An intermediate state, denoted henceforth as the limit condition entails a mixture of fully bonded material and still-porous metal scraps (Fig. 4b). In this state, extruded pieces may exhibit an exterior of solid material while retaining internal sections lacking bonding. Failure to detect this condition promptly could result in significant issues, as extruded pieces may demonstrate considerably reduced structural integrity compared to fully solid samples.
for the production of extruded pieces with a solid external surface but internal sections of unbounded material. Identifying this scenario is essential, as it could lead to significant issues, resulting in a more fragile behavior than a completely solid FSEed sample.

Addressing these challenges requires precise definition of the boundary condition zone and the selection of parameter combinations guaranteed to fall within this range. This determination is facilitated by consulting the diagram presented in Fig. 3, which illustrates the summarized experimental results in terms of bonded and non-bonded specimens. The red cross designates the non-bonding zone, the green dot signifies the bonding zone, and the yellow triangle represents the limit condition zone.

<figure><img src="image_660.png" /><figcaption>Figure 3: Bonding conditions as a function of the process parameters.</figcaption></figure>

From the simulations point of view, the bonding, no bonding and limit conditions are reported in Fig. 4. The scenario where metal scraps are fully bonded (Fig. 4a) contrasts notably with instances where bonding fails to occur before extrusion (Fig. 4c).

An intermediate state, denoted henceforth as the limit condition entails a mixture of fully bonded material and still-porous metal scraps (Fig. 4b). In this state, extruded pieces may exhibit an exterior of solid material while retaining internal sections lacking bonding. Failure to detect this condition promptly could result in significant issues, as extruded pieces may demonstrate considerably reduced structural integrity compared to fully solid samples.
<figure><img src="image_650.png" /><figcaption>Figure 4: Different bonding conditions: (a) Completely bonded with S=800 rpm and F= 1 mm/s, (b) Limit condition with S=800 rpm and F= 2 mm/s and (c) Completely not bonded with S=800 rpm and F= 3 mm/s.</figcaption></figure>

As already explained, the microstructure of aluminum alloy extruded profiles has garnered significant attention within the extrusion industry. For this reason, an additional Fortran routine was implemented in the post-processing phase within DEFORM™ 3D software. This routine, already developed for the traditional extrusion process and using Qform Extrusion software, was adapted to this new technology and this different software. With the revised routine, the possibility to foresee the medium grain dimension independently from the considered aluminum alloy was confirmed. Moreover, the prediction of the final grain dimensions is also possible to be achieved regardless of the size of the initial grains. Indeed, it is sufficient to modify only the parameters relative to the initial dimensions of the grain (also the longitudinal and transverse ones, if the starting material has no equiaxed grains). This allows the model to be adapted to the starting material, be it in the form of chips, powder or flat material by changing only the relevant variables first in the simulation model, mainly the starting density and the aluminum alloy, and the parameters mentioned above in the post-processor prediction model.

In the analyzed case, metal chip was considered as the starting material. Because of the small size and the machining operations that contributed to the formation of this chip, it is possible to think of it as consisting of rather uniform and equiaxial grains. For this reason, only the grain thickness will be plotted. Plotting this information within the routine, the microstructural trend of the extruded sample was evaluated by FSE. At the 5-mm end of the primary die run, a distinct grain refinement can be observed, as can be seen in Fig. 5. The trend of grain evolution and the results obtained were in line with what can be found in the literature. In particular, Buffa et al., obtained comparable results considering the same starting alloy, in the form of chips and the same rotational speed, showing that, from the microstructure point of view, no great differences were shown as a function of the different force applied in the tool thrust phase (Fig. 6), comparable in our case to the different tool descent speeds considered [20].
<figure><img src="image_661.png" /><figcaption>Figure 4: Different bonding conditions: (a) Completely bonded with S=800 rpm and F= 1 mm/s, (b) Limit condition with S=800 rpm and F= 2 mm/s and (c) Completely not bonded with S=800 rpm and F= 3 mm/s.</figcaption></figure>

As already explained, the microstructure of aluminum alloy extruded profiles has garnered significant attention within the extrusion industry. For this reason, an additional Fortran routine was implemented in the post-processing phase within DEFORM™ 3D software. This routine, already developed for the traditional extrusion process and using Qform Extrusion software, was adapted to this new technology and this different software. With the revised routine, the possibility to foresee the medium grain dimension independently from the considered aluminum alloy was confirmed. Moreover, the prediction of the final grain dimensions is also possible to be achieved regardless of the size of the initial grains. Indeed, it is sufficient to modify only the parameters relative to the initial dimensions of the grain (also the longitudinal and transverse ones, if the starting material has no equiaxed grains). This allows the model to be adapted to the starting material, be it in the form of chips, powder or flat material by changing only the relevant variables first in the simulation model, mainly the starting density and the aluminum alloy, and the parameters mentioned above in the post-processor prediction model.

In the analyzed case, metal chip was considered as the starting material. Because of the small size and the machining operations that contributed to the formation of this chip, it is possible to think of it as consisting of rather uniform and equiaxial grains. For this reason, only the grain thickness will be plotted. Plotting this information within the routine, the microstructural trend of the extruded sample was evaluated by FSE. At the 5-mm end of the primary die run, a distinct grain refinement can be observed, as can be seen in Fig. 5. The trend of grain evolution and the results obtained were in line with what can be found in the literature. In particular, Buffa et al., obtained comparable results considering the same starting alloy, in the form of chips and the same rotational speed, showing that, from the microstructure point of view, no great differences were shown as a function of the different force applied in the tool thrust phase (Fig. 6), comparable in our case to the different tool descent speeds considered [20].
<figure><img src="image_651.png" /><figcaption>Figure 5: Grain dimensions in the FSEed sample obtained with S800 rpm and F=1 mm/min.</figcaption></figure>

<figure><img src="image_652.png" /><figcaption>Figure 6: Etched cross sections of the AA6082 FSEed samples obtained with rotational speeds varying between 300 and 500 rpm and applied forces between 15 and 21 kN [20].</figcaption></figure>

The grain sizes obtained, around 15-30 µm in the fully extruded part, are also consistent with what can be found in the literature [21]. Nevertheless, the proposed predictive model was also able to predict the trend of the average grain size within the extruded pieces. This is demonstrated by the images shown in Fig. 7, in which the close-up of the results obtained by applying the developed routine to the proposed simulation model is compared with the result obtained for the Friction Stir Consolidation process, which is a very similar process to the FSE process, except for the presence of the hole in the primary die [22].
<figure><img src="image_662.png" /><figcaption>Figure 5: Grain dimensions in the FSEed sample obtained with S800 rpm and F=1 mm/min.</figcaption></figure>

<figure><img src="image_663.png" /><figcaption>Figure 6: Etched cross sections of the AA6082 FSEed samples obtained with rotational speeds varying between 300 and 500 rpm and applied forces between 15 and 21 kN [20].</figcaption></figure>

The grain sizes obtained, around 15-30 µm in the fully extruded part, are also consistent with what can be found in the literature [21]. Nevertheless, the proposed predictive model was also able to predict the trend of the average grain size within the extruded pieces. This is demonstrated by the images shown in Fig. 7, in which the close-up of the results obtained by applying the developed routine to the proposed simulation model is compared with the result obtained for the Friction Stir Consolidation process, which is a very similar process to the FSE process, except for the presence of the hole in the primary die [22].
<figure><img src="image_653.png" /><figcaption>Figure 7: a) Close-up of the mean grains dimension of the compacted scraps reported in Figure 4. b) Grain characteristics obtained in a typical compacted workpiece obtained with the Friction Stir Consolidation (a very similar process [22]).</figcaption></figure>

## Conclusions

In the present work, a comprehensive understanding of the recrystallization mechanisms involved in the FSE process was carried out together with the FE simulation using DEFORM™ 3D code. The main outcomes of this work are summarized as follows:

* The configuration of the Finite Element model offers an accurate forecast of the FSE physic.

* Piwnik and Plata criterion has demonstrated its effectiveness in accurately predicting the occurrence of bonding phenomena and understanding the impact of process parameters on the quality of produced parts.

* An innovative AA6061 recrystallization model originally developed to investigate the microstructure evolution in conventional extruded profiles proved to be suitable also for the FSE process.

Future developments of this research work will have as a central focus the further experimental validation of the simulative results obtained, with particular attention to the analysis of the effect of process parameters on the bonding conditions of the obtained extruded pieces.

## References

[1] Brand U, Lang M. Green Economy Update Green Economy Update. vol. 1. Cheltenham/Northampton: Edward Elgar; 2015.

[2] Baffari D, Buffa G, Campanella D, Fratini L, Reynolds AP. Process mechanics in Friction Stir Extrusion of magnesium alloys chips through experiments and numerical simulation. J Manuf Process 2017;29:41–9. https://doi.org/10.1016/j.jmapro.2017.07.010

[3] Baffari D, Buffa G, Campanella D, Fratini L. Al-SiC Metal Matrix Composite production through Friction Stir Extrusion of aluminum chips. Procedia Eng., vol. 207, Elsevier B.V.; 2017, p. 419–24. https://doi.org/10.1016/j.proeng.2017.10.798

[4] Manchiraju K. Direct Solid-State Conversion of Recyclable Metals and Alloys. Golden, CO (United States): Southwire Company; 2012. https://doi.org/10.2172/1039705

[5] Izatt RM. Metal Sustainability: Global Challenges, Consequences, and Prospects. Wiley; 2016. https://doi.org/10.1002/9781119009115

[6] Donati L, Reggiani B, Pelaccia R, Negozio M, Di Donato S. Advancements in extrusion and drawing: a review of the contributes by the ESAFORM community. Int J Mater Form 2022;15:1–28. https://doi.org/10.1007/s12289-022-01664-w
<figure><img src="image_664.png" /><figcaption>Figure 7: a) Close-up of the mean grains dimension of the compacted scraps reported in Figure 4. b) Grain characteristics obtained in a typical compacted workpiece obtained with the Friction Stir Consolidation (a very similar process [22]).</figcaption></figure>

## Conclusions

In the present work, a comprehensive understanding of the recrystallization mechanisms involved in the FSE process was carried out together with the FE simulation using DEFORM™ 3D code. The main outcomes of this work are summarized as follows:

* The configuration of the Finite Element model offers an accurate forecast of the FSE physic.

* Piwnik and Plata criterion has demonstrated its effectiveness in accurately predicting the occurrence of bonding phenomena and understanding the impact of process parameters on the quality of produced parts.

* An innovative AA6061 recrystallization model originally developed to investigate the microstructure evolution in conventional extruded profiles proved to be suitable also for the FSE process.

Future developments of this research work will have as a central focus the further experimental validation of the simulative results obtained, with particular attention to the analysis of the effect of process parameters on the bonding conditions of the obtained extruded pieces.

## References

[1] Brand U, Lang M. Green Economy Update Green Economy Update. vol. 1. Cheltenham/Northampton: Edward Elgar; 2015.

[2] Baffari D, Buffa G, Campanella D, Fratini L, Reynolds AP. Process mechanics in Friction Stir Extrusion of magnesium alloys chips through experiments and numerical simulation. J Manuf Process 2017;29:41–9. https://doi.org/10.1016/j.jmapro.2017.07.010

[3] Baffari D, Buffa G, Campanella D, Fratini L. Al-SiC Metal Matrix Composite production through Friction Stir Extrusion of aluminum chips. Procedia Eng., vol. 207, Elsevier B.V.; 2017, p. 419–24. https://doi.org/10.1016/j.proeng.2017.10.798

[4] Manchiraju K. Direct Solid-State Conversion of Recyclable Metals and Alloys. Golden, CO (United States): Southwire Company; 2012. https://doi.org/10.2172/1039705

[5] Izatt RM. Metal Sustainability: Global Challenges, Consequences, and Prospects. Wiley; 2016. https://doi.org/10.1002/9781119009115

[6] Donati L, Reggiani B, Pelaccia R, Negozio M, Di Donato S. Advancements in extrusion and drawing: a review of the contributes by the ESAFORM community. Int J Mater Form 2022;15:1–28. https://doi.org/10.1007/s12289-022-01664-w
[7] Zhang C, Wang C, Zhang Q, Zhao G, Chen L. Influence of extrusion parameters on microstructure, texture, and second-phase particles in an Al-Mg-Si alloy. J Mater Process Technol 2019;270:323-34. https://doi.org/10.1016/j.jmatprotec.2019.03.014

[8] Zhang C, Wang C, Guo R, Zhao G, Chen L, Sun W, et al. Investigation of dynamic recrystallization and modeling of microstructure evolution of an Al-Mg-Si aluminum alloy during high-temperature deformation. J Alloys Compd 2019;773:59-70. https://doi.org/10.1016/j.jallcom.2018.09.263

[9] Zhang T, Li L, Lu S, Li Z, Chen P, Gong H. Static recrystallization kinetics and microstructure evolution of 7055 aluminum alloy. Metall Res Technol 2019;116:120. https://doi.org/10.1051/metal/2018046

[10] Plata M, Piwnik J. Theoretical and experimental analysis of seam weld formation in hot extrusion of aluminum alloys. 7th Int. Alum. Extrus. Technol., 2000, p. 205-11.

[11] Negozio M, Donati L, Pelaccia R, Reggiani B, Di Donato S. Experimental analysis and modeling of the recrystallization behaviour of a AA6060 extruded profile. Mater. Res. Proc., vol. 28, Association of American Publishers; 2023, p. 477-86. https://doi.org/10.21741/9781644902479-52

[12] Bocchi S, D'urso G, Giardini C, Maccarini G. A Simulative Method for Studying the Bonding Condition of Friction Stir Extrusion. Key Eng Mater 2022;926 KEM:2333-41. https://doi.org/10.4028/P-FT5355

[13] Bocchi S, D'Urso G, Giardini C. Simulative Model for the Feasibility Study and Stress Analysis of Full Dense Rods and Pipes Produced by Friction Stir Extrusion. Lect. Notes Mech. Eng., 2024, p. 398-406. https://doi.org/10.1007/978-3-031-41023-9_41

[14] Bocchi S, D'Urso G, Giardini C. Numerical Modeling of a Sustainable Solid-State Recycling of Aluminum Scraps by Means of Friction Stir Extrusion Process. Materials (Basel) 2023;16:4375. https://doi.org/10.3390/ma16124375

[15] Negozio M, Pelaccia R, Donati L, Reggiani B. Simulation of the microstructure evolution during the extrusion of two industrial-scale AA6063 profiles. J Manuf Process 2023;99:501-12. https://doi.org/10.1016/j.jmapro.2023.05.075

[16] Negozio M, Pelaccia R, Donati L, Reggiani B. Numerical investigation of the surface recrystallization during the extrusion of a AA6082 aluminum alloy under different process conditions. Int J Adv Manuf Technol 2023;129:1585-99. https://doi.org/10.1007/s00170-023-12397-8

[17] Khamei AA, Dehghani K, Mahmudi R. Modeling the hot ductility of AA6061 aluminum alloy after severe plastic deformation. JOM 2015;67:966-72. https://doi.org/10.1007/s11837-015-1354-3

[18] Vatne HE, Furu T, Ørsund R, Nes E. Modelling recrystallization after hot deformation of aluminium. Acta Mater 1996;44:4463-73. https://doi.org/10.1016/1359-6454(96)00078-X.

[19] Donati L, Segatori A, El Mehtedi M, Tomesani L. Grain evolution analysis and experimental validation in the extrusion of 6XXX alloys by use of a lagrangian FE code. Int J Plast 2013;46:70-81. https://doi.org/10.1016/j.ijplas.2012.11.008

[20] Buffa G, Campanella D, Adnan M, La Commare U, Ingarao G, Fratini L. Improving the Industrial Efficiency of Recycling Aluminum Alloy Chips Using Friction Stir Extrusion: Thin
[7] Zhang C, Wang C, Zhang Q, Zhao G, Chen L. Influence of extrusion parameters on microstructure, texture, and second-phase particles in an Al-Mg-Si alloy. J Mater Process Technol 2019;270:323-34. https://doi.org/10.1016/j.jmatprotec.2019.03.014

[8] Zhang C, Wang C, Guo R, Zhao G, Chen L, Sun W, et al. Investigation of dynamic recrystallization and modeling of microstructure evolution of an Al-Mg-Si aluminum alloy during high-temperature deformation. J Alloys Compd 2019;773:59-70. https://doi.org/10.1016/j.jallcom.2018.09.263

[9] Zhang T, Li L, Lu S, Li Z, Chen P, Gong H. Static recrystallization kinetics and microstructure evolution of 7055 aluminum alloy. Metall Res Technol 2019;116:120. https://doi.org/10.1051/metal/2018046

[10] Plata M, Piwnik J. Theoretical and experimental analysis of seam weld formation in hot extrusion of aluminum alloys. 7th Int. Alum. Extrus. Technol., 2000, p. 205-11.

[11] Negozio M, Donati L, Pelaccia R, Reggiani B, Di Donato S. Experimental analysis and modeling of the recrystallization behaviour of a AA6060 extruded profile. Mater. Res. Proc., vol. 28, Association of American Publishers; 2023, p. 477-86. https://doi.org/10.21741/9781644902479-52

[12] Bocchi S, D'urso G, Giardini C, Maccarini G. A Simulative Method for Studying the Bonding Condition of Friction Stir Extrusion. Key Eng Mater 2022;926 KEM:2333-41. https://doi.org/10.4028/P-FT5355

[13] Bocchi S, D'Urso G, Giardini C. Simulative Model for the Feasibility Study and Stress Analysis of Full Dense Rods and Pipes Produced by Friction Stir Extrusion. Lect. Notes Mech. Eng., 2024, p. 398-406. https://doi.org/10.1007/978-3-031-41023-9_41

[14] Bocchi S, D'Urso G, Giardini C. Numerical Modeling of a Sustainable Solid-State Recycling of Aluminum Scraps by Means of Friction Stir Extrusion Process. Materials (Basel) 2023;16:4375. https://doi.org/10.3390/ma16124375

[15] Negozio M, Pelaccia R, Donati L, Reggiani B. Simulation of the microstructure evolution during the extrusion of two industrial-scale AA6063 profiles. J Manuf Process 2023;99:501-12. https://doi.org/10.1016/j.jmapro.2023.05.075

[16] Negozio M, Pelaccia R, Donati L, Reggiani B. Numerical investigation of the surface recrystallization during the extrusion of a AA6082 aluminum alloy under different process conditions. Int J Adv Manuf Technol 2023;129:1585-99. https://doi.org/10.1007/s00170-023-12397-8

[17] Khamei AA, Dehghani K, Mahmudi R. Modeling the hot ductility of AA6061 aluminum alloy after severe plastic deformation. JOM 2015;67:966-72. https://doi.org/10.1007/s11837-015-1354-3

[18] Vatne HE, Furu T, Ørsund R, Nes E. Modelling recrystallization after hot deformation of aluminium. Acta Mater 1996;44:4463-73. https://doi.org/10.1016/1359-6454(96)00078-X.

[19] Donati L, Segatori A, El Mehtedi M, Tomesani L. Grain evolution analysis and experimental validation in the extrusion of 6XXX alloys by use of a lagrangian FE code. Int J Plast 2013;46:70-81. https://doi.org/10.1016/j.ijplas.2012.11.008

[20] Buffa G, Campanella D, Adnan M, La Commare U, Ingarao G, Fratini L. Improving the Industrial Efficiency of Recycling Aluminum Alloy Chips Using Friction Stir Extrusion: Thin
Wires Production Process. Int J Precis Eng Manuf Technol 2024. https://doi.org/10.1007/s40684-023-00573-w

[21] Chan CYC, Rath L, Suhuddin UFH, Klusemann B. Friction extrusion processing of aluminum powders: Microstructure homogeneity and mechanical properties. Mater. Res. Proc., vol. 28, Association of American Publishers; 2023, p. 515-22. https://doi.org/10.21741/9781644902479-56

[22] Latif A, Gucciardi M, Ingarao G, Fratini L. Outlining the Limits of Friction Stir Consolidation as Used as an Aluminum Alloys Recycling Approach. Smart Innov. Syst. Technol., vol. 262 SIST, Springer Science and Business Media Deutschland GmbH; 2022, p. 169-80. https://doi.org/10.1007/978-981-16-6128-0_17
Wires Production Process. Int J Precis Eng Manuf Technol 2024. https://doi.org/10.1007/s40684-023-00573-w

[21] Chan CYC, Rath L, Suhuddin UFH, Klusemann B. Friction extrusion processing of aluminum powders: Microstructure homogeneity and mechanical properties. Mater. Res. Proc., vol. 28, Association of American Publishers; 2023, p. 515–22. https://doi.org/10.21741/9781644902479-56

[22] Latif A, Gucciardi M, Ingarao G, Fratini L. Outlining the Limits of Friction Stir Consolidation as Used as an Aluminum Alloys Recycling Approach. Smart Innov. Syst. Technol., vol. 262 SIST, Springer Science and Business Media Deutschland GmbH; 2022, p. 169–80. https://doi.org/10.1007/978-981-16-6128-0_17
# Mechanical behavior and flow curve modeling of AZ31B magnesium alloy under low frequency vibration assisted tension

TIAN Ye¹,ª, WEN Zhang¹,ᵇ, SHEN Weicai¹,ᶜ, ZHUANG Xincun¹,ᵈ*, ZHAO Zhen¹,ᵉ*

¹Institute of Forming Technology and Equipment, NERC of Die & Mold CAD, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200030, China

ªsjtutianye@sjtu.edu.cn, ᵇw_zhang@sjtu.edu.cn, ᶜweicaishen@sjtu.edu.cn

ᵈgeorgezxc@sjtu.edu.cn, ᶜzzhao@sjtu.edu.cn

**Keywords:** Low Frequency Vibration, Mechanical Behavior, Vibration Hardening, Flow Curve Model

**Abstract.** This study aimed to investigate the impact of the low-frequency vibration field on the mechanical behavior of AZ31B magnesium alloy. To achieve this, a series of low-frequency vibration-assisted tensile tests were conducted at room temperature. The effect of various vibration parameters on the mechanical characteristics of AZ31B was studied. The findings indicate that the introduction of low-frequency vibration field during the traditional tensile process leads to periodic oscillations in the flow stress of the material. Moreover, the maximum stress value was greater than the condition without vibration. The vibration hardening value is not significantly affected by the vibration frequency, but rather depends on the amplitude. The magnitude of the vibration hardening value increases with the increasing amplitude, and this relationship is non-linear. On this basis, a modified Hockett-Sherby hardening model was constructed by incorporating the vibration parameters. Statistical analysis of stress-strain curve, vibration hardening and stress fluctuation values were conducted to identify the model parameters. By comparing with the experimental result, the stress change of AZ31B magnesium alloy under low frequency vibration tension could be well predicted by this model.

**Introduction**

Vibration-assisted plastic forming technology is a new plastic forming method developed in vibration utilization engineering in recent years [1]. It utilizes the stress field and vibration energy field generated by vibration to reduce the load, increase the forming limit and improve the forming quality of the parts [2]. Cao et al. [3] studied 6061 aluminum alloy and found that after applying ultrasonic vibration, the evolution of dislocations was positively regulated, resulting in a decrease in deformation resistance. The level of stress reduction was proportional to the amplitude of ultrasonic waves. Ultrasonic vibration softening was mainly caused by stress superposition and acoustic softening effect. Han et al. [4] proposed a new tool-workpiece combined vibration ultrasonic micro-extrusion process, which can reduce the micro-extrusion stress, extend the extrusion length, enhance the grain deformation effect and generate fiber streamline. Deng [5] et al. applied low-frequency vibration field in T2 pure copper compression and found that low-frequency vibration increased the dislocation density in the material and promoted the formation of dislocation cells. The dislocation cells gradually evolved into new grain boundaries, and the grains were refined. Lin et al. [6] conducted 10 Hz and 0.3 mm amplitude low-frequency vibration assisted compression tests on Ti45Nb titanium alloy. The low-frequency vibration field promoted the dislocation motion and grain rotation, making the material's deformation more uniform and reducing the probability of crack nucleation.

In order to deeply understand the impact of vibration field on the mechanical behavior of materials and describe the mechanical response of materials under vibration conditions, it is necessary to establish a proper constitutive equation. At present, researchers have conducted

<figure><img src="image_654.png" /></figure>
# Mechanical behavior and flow curve modeling of AZ31B magnesium alloy under low frequency vibration assisted tension

TIAN Ye¹,ª, WEN Zhang¹,ᵇ, SHEN Weicai¹,ᶜ, ZHUANG Xincun¹,ᵈ*, ZHAO Zhen¹,ᵉ*

¹Institute of Forming Technology and Equipment, NERC of Die & Mold CAD, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200030, China

ªsjtutianye@sjtu.edu.cn, ᵇw_zhang@sjtu.edu.cn, ᶜweicaishen@sjtu.edu.cn

ᵈgeorgezxc@sjtu.edu.cn, ᶜzzhao@sjtu.edu.cn

**Keywords:** Low Frequency Vibration, Mechanical Behavior, Vibration Hardening, Flow Curve Model

**Abstract.** This study aimed to investigate the impact of the low-frequency vibration field on the mechanical behavior of AZ31B magnesium alloy. To achieve this, a series of low-frequency vibration-assisted tensile tests were conducted at room temperature. The effect of various vibration parameters on the mechanical characteristics of AZ31B was studied. The findings indicate that the introduction of low-frequency vibration field during the traditional tensile process leads to periodic oscillations in the flow stress of the material. Moreover, the maximum stress value was greater than the condition without vibration. The vibration hardening value is not significantly affected by the vibration frequency, but rather depends on the amplitude. The magnitude of the vibration hardening value increases with the increasing amplitude, and this relationship is non-linear. On this basis, a modified Hockett-Sherby hardening model was constructed by incorporating the vibration parameters. Statistical analysis of stress-strain curve, vibration hardening and stress fluctuation values were conducted to identify the model parameters. By comparing with the experimental result, the stress change of AZ31B magnesium alloy under low frequency vibration tension could be well predicted by this model.

**Introduction**

Vibration-assisted plastic forming technology is a new plastic forming method developed in vibration utilization engineering in recent years [1]. It utilizes the stress field and vibration energy field generated by vibration to reduce the load, increase the forming limit and improve the forming quality of the parts [2]. Cao et al. [3] studied 6061 aluminum alloy and found that after applying ultrasonic vibration, the evolution of dislocations was positively regulated, resulting in a decrease in deformation resistance. The level of stress reduction was proportional to the amplitude of ultrasonic waves. Ultrasonic vibration softening was mainly caused by stress superposition and acoustic softening effect. Han et al. [4] proposed a new tool-workpiece combined vibration ultrasonic micro-extrusion process, which can reduce the micro-extrusion stress, extend the extrusion length, enhance the grain deformation effect and generate fiber streamline. Deng [5] et al. applied low-frequency vibration field in T2 pure copper compression and found that low-frequency vibration increased the dislocation density in the material and promoted the formation of dislocation cells. The dislocation cells gradually evolved into new grain boundaries, and the grains were refined. Lin et al. [6] conducted 10 Hz and 0.3 mm amplitude low-frequency vibration assisted compression tests on Ti45Nb titanium alloy. The low-frequency vibration field promoted the dislocation motion and grain rotation, making the material's deformation more uniform and reducing the probability of crack nucleation.

In order to deeply understand the impact of vibration field on the mechanical behavior of materials and describe the mechanical response of materials under vibration conditions, it is necessary to establish a proper constitutive equation. At present, researchers have conducted

<figure><img src="image_665.png" /></figure>
extensive research on the construction of models for vibration plastic forming. Meng et al. [7] established a hybrid constitutive model based on dislocation evolution theory, which can respectively describe the stress superposition and acoustic softening effects under the action of ultrasonic vibration field. Based on the theory of crystal plasticity, thermal activation and dislocation evolution, Yao et al. [8] proposed an acoustic model considering softening and residual effects, and calibrated it by ultrasonic vibration-assisted compression test. Through the study of 6063 aluminum alloy, Xie et al. [9] established a constitutive model for ultrasonic vibration compression based on the Johnson-Cook model. Prabhakar et al. [10] extended the constitutive model based on dislocation density to the softening model induced by ultrasonic effect by superimposing transverse ultrasonic vibration in the tensile experiment of aluminum alloy. The predicted results are in good agreement with the experimental results. Zhang et al. [11] introduced the mechanical work done by low-frequency vibration, the critical vibration energy value and Hooke's law into the framework of thermal activation to establish the physical intrinsic structure from the perspective of thermal activation.

To investigate the effect of low-frequency vibration field on mechanical behavior of AZ31B magnesium alloy, this study conducted a series of room temperature low-frequency vibration assisted tensile tests. The investigation focused on analyzing the impact of various vibration parameters on the mechanical behavior of AZ31B. Subsequently, a flow curve model incorporating the vibration parameters was developed. A modified Hockett-Sherby hardening model was constructed by statistical analysis of stress-strain curve, vibration hardening value, stress fluctuation value, etc.

## Material and experimental setup

The material used in this study is commercial AZ31B-H24 magnesium alloy sheet, and its chemical composition is shown in Table 1. The thickness of the as-received sheet is 2 mm. The tensile specimens were designed according to ASTM standards and prepared along the rolling direction of the as-received sheet. The specific dimensions are shown in Fig. 1a.

The experimental platform mainly consists of an electronic material testing machine (UTM5105) and an electro-hydraulic servo excitation device, which can accurately control low-frequency vibration within the frequency range of 0-50 Hz and amplitude range of 0-1 mm. The specific device is shown in Fig. 1b. The hydraulic control unit can provide a maximum excitation force of up to 200 kN. In order to ensure the reliability of data during the vibration process, the collection frequency of the electronic testing machine was set to 330 Hz. The strain of the material was measured using the digital image correlation (DIC) method. All test images were recorded with a pair of cameras from GOM's ARAMIS optical strain measurement system. The resolution of the cameras was 2752 × 2200 pixels and the system was calibrated to have a reference area of 100 × 100 mm. Considering that the superposition of low-frequency vibration will only make the material produce small elastic deformation, but will not change the deformation mode and stress state of the material. The acquisition frequency of three frames per second was used in this study, which can accurately record the strain condition during the tensile process to meet the requirements of strain testing.

The tensile speed in the quasi-static tensile test was 2 mm/min, with a strain rate of approximately $5 \times 10^{-4}/s$. In order to improve the accuracy of the experiment, the entrance force of the testing machine was set to 50 N, and after each experiment, the crossbeam of the testing machine automatically resets to the position before the last test started. The vibration device was activated when the material was stretched to 5% strain for 60 s, and finally stopped at 10% strain. The amplitude range used in this experiment was 0-0.1 mm and the frequency range was 0-30 Hz. The specific testing conditions are shown in Table 2, and each experiment should be tested at least three times to ensure the reproducibility of the experiment.
extensive research on the construction of models for vibration plastic forming. Meng et al. [7] established a hybrid constitutive model based on dislocation evolution theory, which can respectively describe the stress superposition and acoustic softening effects under the action of ultrasonic vibration field. Based on the theory of crystal plasticity, thermal activation and dislocation evolution, Yao et al. [8] proposed an acoustic model considering softening and residual effects, and calibrated it by ultrasonic vibration-assisted compression test. Through the study of 6063 aluminum alloy, Xie et al. [9] established a constitutive model for ultrasonic vibration compression based on the Johnson-Cook model. Prabhakar et al. [10] extended the constitutive model based on dislocation density to the softening model induced by ultrasonic effect by superimposing transverse ultrasonic vibration in the tensile experiment of aluminum alloy. The predicted results are in good agreement with the experimental results. Zhang et al. [11] introduced the mechanical work done by low-frequency vibration, the critical vibration energy value and Hooke's law into the framework of thermal activation to establish the physical intrinsic structure from the perspective of thermal activation.

To investigate the effect of low-frequency vibration field on mechanical behavior of AZ31B magnesium alloy, this study conducted a series of room temperature low-frequency vibration assisted tensile tests. The investigation focused on analyzing the impact of various vibration parameters on the mechanical behavior of AZ31B. Subsequently, a flow curve model incorporating the vibration parameters was developed. A modified Hockett-Sherby hardening model was constructed by statistical analysis of stress-strain curve, vibration hardening value, stress fluctuation value, etc.

## Material and experimental setup

The material used in this study is commercial AZ31B-H24 magnesium alloy sheet, and its chemical composition is shown in Table 1. The thickness of the as-received sheet is 2 mm. The tensile specimens were designed according to ASTM standards and prepared along the rolling direction of the as-received sheet. The specific dimensions are shown in Fig. 1a.

The experimental platform mainly consists of an electronic material testing machine (UTM5105) and an electro-hydraulic servo excitation device, which can accurately control low-frequency vibration within the frequency range of 0-50 Hz and amplitude range of 0-1 mm. The specific device is shown in Fig. 1b. The hydraulic control unit can provide a maximum excitation force of up to 200 kN. In order to ensure the reliability of data during the vibration process, the collection frequency of the electronic testing machine was set to 330 Hz. The strain of the material was measured using the digital image correlation (DIC) method. All test images were recorded with a pair of cameras from GOM's ARAMIS optical strain measurement system. The resolution of the cameras was 2752 × 2200 pixels and the system was calibrated to have a reference area of 100 × 100 mm. Considering that the superposition of low-frequency vibration will only make the material produce small elastic deformation, but will not change the deformation mode and stress state of the material. The acquisition frequency of three frames per second was used in this study, which can accurately record the strain condition during the tensile process to meet the requirements of strain testing.

The tensile speed in the quasi-static tensile test was 2 mm/min, with a strain rate of approximately $5 \times 10^{-4}/s$. In order to improve the accuracy of the experiment, the entrance force of the testing machine was set to 50 N, and after each experiment, the crossbeam of the testing machine automatically resets to the position before the last test started. The vibration device was activated when the material was stretched to 5% strain for 60 s, and finally stopped at 10% strain. The amplitude range used in this experiment was 0-0.1 mm and the frequency range was 0-30 Hz. The specific testing conditions are shown in Table 2, and each experiment should be tested at least three times to ensure the reproducibility of the experiment.
<table><caption>Table 1. Chemical composition of As-received AZ31B-H24 (in wt%).</caption><thead><tr><th>Al</th><th>Zn</th><th>Mn</th><th>Si</th><th>Fe</th><th>Cu</th><th>Ni</th></tr></thead><tbody><tr><td>3.10</td><td>0.91</td><td>0.31</td><td>0.02</td><td>0.00</td><td>0.00</td><td>0.00</td></tr></tbody></table>

<figure><img src="image_655.png" /><figcaption>Fig. 1 Schematics of the tension specimen (a) and experimental setup (b).</figcaption></figure>

<table><caption>Table 2. Conditions of the LFVT tests.</caption><thead><tr><th>Item</th><th>Condition</th></tr></thead><tbody><tr><td>Material</td><td>AZ31B-H24</td></tr><tr><td>Effect of amplitude</td><td>0/0.025/0.05/0.075/0.1 mm, 20 Hz</td></tr><tr><td>Effect of frequency</td><td>0.075 mm, 0/10/20/30 Hz</td></tr></tbody></table>

## Mechanical behavior under LFVT

Fig. 2a shows a comparison of the curves with and without vibration, and it can be found that the two curves basically overlap before the superimposed low-frequency vibration. Once low-frequency vibration was applied, the stress exhibits periodic fluctuations, which is called the stress superposition effect [12]. The stress superposition effect was first proposed by Nevill and Brotzen [13], who suggested that the average load drop in ultrasonic vibration is attributed to the superposition mechanism of steady and alternating stresses. Relevant researchers believe [14,15] that the effect is also applicable to low-frequency vibration, and the metal under the low-frequency field will also produce a nonlinear response of stress caused by loading and unloading, which is directly related to the elasticity of the material. Deng et al. [5] elucidated the stress superposition effect from the perspective of the stress wave, and the fluctuating stress field will be generated inside the metal material under the low-frequency vibration field, and the effect will certainly exist in the in vibroforming. In the uniaxial tensile process, the low-frequency vibration field produces periodic stress fluctuations in the stress component along the tensile direction, but does not cause changes in the plastic deformation mechanism or internal microstructure of the material.

From the above, it can be seen that during a single oscillation period, there is an upper limit denoted as $\sigma_{max}$ and a lower limit denoted as $\sigma_{min}$ for the flow stress. The disparity between these two values is defined as the stress fluctuation, which can be mathematically expressed as $\Delta\sigma = $
<table><caption>Table 1. Chemical composition of As-received AZ31B-H24 (in wt%).</caption><thead><tr><th>Al</th><th>Zn</th><th>Mn</th><th>Si</th><th>Fe</th><th>Cu</th><th>Ni</th></tr></thead><tbody><tr><td>3.10</td><td>0.91</td><td>0.31</td><td>0.02</td><td>0.00</td><td>0.00</td><td>0.00</td></tr></tbody></table>

<figure><img src="image_666.png" /><figcaption>Fig. 1 Schematics of the tension specimen (a) and experimental setup (b).</figcaption></figure>

<table><caption>Table 2. Conditions of the LFVT tests.</caption><thead><tr><th>Item</th><th>Condition</th></tr></thead><tbody><tr><td>Material</td><td>AZ31B-H24</td></tr><tr><td>Effect of amplitude</td><td>0/0.025/0.05/0.075/0.1 mm, 20 Hz</td></tr><tr><td>Effect of frequency</td><td>0.075 mm, 0/10/20/30 Hz</td></tr></tbody></table>

## Mechanical behavior under LFVT

Fig. 2a shows a comparison of the curves with and without vibration, and it can be found that the two curves basically overlap before the superimposed low-frequency vibration. Once low-frequency vibration was applied, the stress exhibits periodic fluctuations, which is called the stress superposition effect [12]. The stress superposition effect was first proposed by Nevill and Brotzen [13], who suggested that the average load drop in ultrasonic vibration is attributed to the superposition mechanism of steady and alternating stresses. Relevant researchers believe [14,15] that the effect is also applicable to low-frequency vibration, and the metal under the low-frequency field will also produce a nonlinear response of stress caused by loading and unloading, which is directly related to the elasticity of the material. Deng et al. [5] elucidated the stress superposition effect from the perspective of the stress wave, and the fluctuating stress field will be generated inside the metal material under the low-frequency vibration field, and the effect will certainly exist in the in vibroforming. In the uniaxial tensile process, the low-frequency vibration field produces periodic stress fluctuations in the stress component along the tensile direction, but does not cause changes in the plastic deformation mechanism or internal microstructure of the material.

From the above, it can be seen that during a single oscillation period, there is an upper limit denoted as $\sigma_{max}$ and a lower limit denoted as $\sigma_{min}$ for the flow stress. The disparity between these two values is defined as the stress fluctuation, which can be mathematically expressed as $\Delta\sigma = $
$\sigma_{max} - \sigma_{min}$. The effects of amplitude and frequency on $\Delta\sigma$ are depicted in Fig. 2b and c, respectively. It is clear from the data that $\Delta\sigma$ is positively correlated with the increase in amplitude. However, changing the vibration frequency does not have a substantial effect on $\Delta\sigma$.

Furthermore, it has been observed from Fig. 2a that the peak of the oscillation curve exceeded the traditional tensile flow stress, indicating the presence of the vibration hardening phenomenon in the LFVT process. To quantitatively assess the impact of vibration hardening, a metric termed the vibration hardening value $S$ was introduced, which is represented as the discrepancy between the maximum value ($\sigma_{max}$) of the oscillation curve and the flow stress ($\sigma_{WV}$) corresponding to that particular point in without vibration tension, that is, $S = \sigma_{max} - \sigma_{WV}$. The statistical results of $S$ are presented in Table 3. It is noteworthy that the vibration frequency does not exert a direct impact on the vibration hardening value $S$. Similar to $\Delta\sigma$, the vibration hardening value $S$ is primarily influenced by the magnitude of the amplitude. This means that as the amplitude increases, the vibration hardening value $S$ also increases. The amplitude plays a crucial role in determining the magnitude of $S$.

To sum up, whether it is for $\Delta\sigma$ or $S$, the impact of vibration frequency is much smaller than the amplitude. Similar conclusions have been drawn by other researchers regarding the influence of vibration frequency on mechanical behavior [5, 16].

<figure><img src="image_656.png" /><figcaption>Fig. 2 True stress-strain curves under LFVT(a), effect of amplitude on stress fluctuation values(b), effect of frequency on stress fluctuation values(c).</figcaption></figure>

<table><caption>Table 3. Average vibration hardening value $S$</caption><thead><tr><th></th><th colspan="4">Effect of amplitude</th><th colspan="3">Effect of frequency</th></tr></thead><tbody><tr><th>Amplitude(mm)</th><td>0.025</td><td>0.05</td><td>0.075</td><td>0.1</td><td colspan="3">0.075</td></tr><tr><th>Frequency (Hz)</th><td colspan="4">20</td><td>10</td><td>20</td><td>30</td></tr><tr><th>S(MPa)</th><td>5.67</td><td>9.09</td><td>11.38</td><td>13.40</td><td>11.21</td><td>11.38</td><td>11.16</td></tr></tbody></table>

### Flow curve modeling of mechanical behavior under LFVT

The Hockett-Sherby hardening model is a type of saturation model, meaning that as the strain increases, the flow stress reaches a constant value. The original expression for the Hockett-Sherby hardening model is as follows:

$$ \sigma = \sigma_s - (\sigma_s - \sigma_i) \times \exp(-a\epsilon^n) \quad (1) $$

In equation (1), $\sigma_s$ is the saturation stress; $\sigma_i$ is the initial yield stress; $a$ and $n$ are constants ($a > 0$). The parameters of the model ($\sigma_s$, $\sigma_i$, $a$ and $n$) can be determined by applying the results of static tensile tests to the Hockett-Sherby hardening model.

$$ \sigma = 336.14 - 123.7 \times \exp(-12.09 \times \epsilon^{0.87}) \quad (2) $$
$\sigma_{max} - \sigma_{min}$. The effects of amplitude and frequency on $\Delta\sigma$ are depicted in Fig. 2b and c, respectively. It is clear from the data that $\Delta\sigma$ is positively correlated with the increase in amplitude. However, changing the vibration frequency does not have a substantial effect on $\Delta\sigma$.

Furthermore, it has been observed from Fig. 2a that the peak of the oscillation curve exceeded the traditional tensile flow stress, indicating the presence of the vibration hardening phenomenon in the LFVT process. To quantitatively assess the impact of vibration hardening, a metric termed the vibration hardening value $S$ was introduced, which is represented as the discrepancy between the maximum value ($\sigma_{max}$) of the oscillation curve and the flow stress ($\sigma_{WV}$) corresponding to that particular point in without vibration tension, that is, $S = \sigma_{max} - \sigma_{WV}$. The statistical results of $S$ are presented in Table 3. It is noteworthy that the vibration frequency does not exert a direct impact on the vibration hardening value $S$. Similar to $\Delta\sigma$, the vibration hardening value $S$ is primarily influenced by the magnitude of the amplitude. This means that as the amplitude increases, the vibration hardening value $S$ also increases. The amplitude plays a crucial role in determining the magnitude of $S$.

To sum up, whether it is for $\Delta\sigma$ or $S$, the impact of vibration frequency is much smaller than the amplitude. Similar conclusions have been drawn by other researchers regarding the influence of vibration frequency on mechanical behavior [5, 16].

<figure><img src="image_667.png" /><figcaption>Fig. 2 True stress-strain curves under LFVT(a), effect of amplitude on stress fluctuation values(b), effect of frequency on stress fluctuation values(c).</figcaption></figure>

<table><caption>Table 3. Average vibration hardening value $S$</caption><thead><tr><th></th><th colspan="4">Effect of amplitude</th><th colspan="3">Effect of frequency</th></tr></thead><tbody><tr><th>Amplitude(mm)</th><td>0.025</td><td>0.05</td><td>0.075</td><td>0.1</td><td colspan="3">0.075</td></tr><tr><th>Frequency (Hz)</th><td colspan="4">20</td><td>10</td><td>20</td><td>30</td></tr><tr><th>S(MPa)</th><td>5.67</td><td>9.09</td><td>11.38</td><td>13.40</td><td>11.21</td><td>11.38</td><td>11.16</td></tr></tbody></table>

### Flow curve modeling of mechanical behavior under LFVT

The Hockett-Sherby hardening model is a type of saturation model, meaning that as the strain increases, the flow stress reaches a constant value. The original expression for the Hockett-Sherby hardening model is as follows:

$$ \sigma = \sigma_s - (\sigma_s - \sigma_i) \times \exp(-a\varepsilon^n) \quad (1) $$

In equation (1), $\sigma_s$ is the saturation stress; $\sigma_i$ is the initial yield stress; $a$ and $n$ are constants ($a > 0$). The parameters of the model ($\sigma_s$, $\sigma_i$, $a$ and $n$) can be determined by applying the results of static tensile tests to the Hockett-Sherby hardening model.

$$ \sigma = 336.14 - 123.7 \times \exp(-12.09 \times \varepsilon^{0.87}) \quad (2) $$
<figure><img src="image_657.png" /><figcaption>Fig.3 Fitting results of flow stress using the Hockett-Sherby model</figcaption></figure>

When low-frequency vibration was applied, the flow stress of the material undergoes oscillations between the maximum and minimum values. This stress fluctuation, denoted as Δσ, was primarily caused by the elastic loading and unloading of the vibration. According to Hooke's law, it can be inferred that the expression for Δσ is as follows:

$$ \Delta\sigma = \sigma_{max} - \sigma_{min} = c \times \frac{EA}{L} \qquad (3) $$

Here, $\frac{EA}{L}$ represents the variation in strain induced by the applied vibration; c is constant; E is elastic modulus (for magnesium alloys, the value is 45 GPa); L is specimen gauge length (50 mm). Based on the results shown in Fig. 2b, a Δσ - EA/L relationship can be constructed as shown in Fig. 4. The slope of the fitted curve represents the coefficient c. Therefore, the specific expression for the Δσ is given by:

$$ \Delta\sigma = 981 \times A \qquad (4) $$

<figure><img src="image_658.png" /><figcaption>Fig. 4 Linear fitting results of Δσ - EA/L during LFVT</figcaption></figure>

To further understand the relationship between the vibration hardening value S and the amplitude A, a vibration hardening coefficient function f(A) was constructed. The equation (5) represents the relationship between S and f(A), and is given as follows:

$$ S = \sigma_{max} - \sigma_{wv} = \sigma_{wv} \times f(A) - \sigma_{wv} \qquad (5) $$
<figure><img src="image_668.png" /><figcaption>Fig.3 Fitting results of flow stress using the Hockett-Sherby model</figcaption></figure>

When low-frequency vibration was applied, the flow stress of the material undergoes oscillations between the maximum and minimum values. This stress fluctuation, denoted as Δσ, was primarily caused by the elastic loading and unloading of the vibration. According to Hooke's law, it can be inferred that the expression for Δσ is as follows:

$$ \Delta\sigma = \sigma_{max} - \sigma_{min} = c \times \frac{EA}{L} \qquad (3) $$

Here, $\frac{EA}{L}$ represents the variation in strain induced by the applied vibration; c is constant; E is elastic modulus (for magnesium alloys, the value is 45 GPa); L is specimen gauge length (50 mm). Based on the results shown in Fig. 2b, a $\Delta\sigma - EA/L$ relationship can be constructed as shown in Fig. 4. The slope of the fitted curve represents the coefficient c. Therefore, the specific expression for the $\Delta\sigma$ is given by:

$$ \Delta\sigma = 981 \times A \qquad (4) $$

<figure><img src="image_669.png" /><figcaption>Fig. 4 Linear fitting results of $\Delta\sigma - EA/L$ during LFVT</figcaption></figure>

To further understand the relationship between the vibration hardening value S and the amplitude A, a vibration hardening coefficient function f(A) was constructed. The equation (5) represents the relationship between S and f(A), and is given as follows:

$$ S = \sigma_{max} - \sigma_{wv} = \sigma_{wv} \times f(A) - \sigma_{wv} \qquad (5) $$
Due to the magnitude of *S* is determined by the amplitude *A*. By varying the amplitude, we can observe the corresponding changes in the vibration hardening value. Therefore, it can be deduced that *f*(*A*) also depends on the amplitude *A*. The specific expression for *f*(*A*) is as follows:

$$f(A) = \left\{ k \times \left( \frac{A}{A_0} \right)^q + 1 \right\}^m \qquad (6)$$

where *A* denotes the amplitude; *A*₀ refers to reference amplitude (whose value is 1mm); *k* is constants; *m* and *q* are constants. Obviously, when *A* = 0, *f* (0) = 1. The σmax/σwv - A/A₀ relationship was established and shown in Fig. 5. According to the fitting results, a mathematical expression about *f*(*A*) was obtained:

$$f(A) = (2.27 \times A^{0.66} + 1)^{0.11} \qquad (7)$$

<figure><img src="image_659.png" /><figcaption>Fig. 5 Linear fitting results of σmax/σwv-A/A₀</figcaption></figure>

The fitting results of the modified Hockett-Sherby hardening model parameters are summarized in Table 4.

<table><caption>Table 4. Parameters of modified Hockett-Sherby model</caption><thead><tr><th>parameters</th><th>σ<sub>s</sub>(MPa)</th><th>σ<sub>i</sub>(MPa)</th><th>a</th><th>n</th><th>k</th><th>q</th></tr></thead><tbody><tr><td>value</td><td>336.14</td><td>212.44</td><td>12.09</td><td>0.87</td><td>2.27</td><td>0.66</td></tr><tr><td>parameters</td><td>m</td><td>c</td><td>E(GPa)</td><td>L(mm)</td><td>A<sub>0</sub>(mm)</td><td></td></tr><tr><td>value</td><td>0.107</td><td>1.09</td><td>45</td><td>50</td><td>1</td><td></td></tr></tbody></table>

In summary, when low-frequency axial vibration was applied to AZ31B magnesium alloy, the curve of stress variation with strain of the material can be described by equation (8):

$$ \begin{cases} \sigma_{max} = \{336.14 - 123.7 \times \exp(-12.09 \times \varepsilon^{0.87})\} \times (2.27 \times A^{0.66} + 1)^{0.11} \\ \sigma_{min} = 336.14 - 123.7 \times \exp(-12.09 \times \varepsilon^{0.87}) \times (2.27 \times A^{0.66} + 1)^{0.11} - \Delta\sigma \\ \Delta\sigma = 981 \times A \end{cases} \qquad (8) $$

To verify the accuracy of the model, the LFVT results with an amplitude of 0.06 mm and a frequency of 20 Hz were compared with the predicted results. The comparison result is shown in Fig. 6. The results demonstrated that the model accurately captures the mechanical behavior of AZ31B under LFVT. This suggests that the Hockett-Sherby hardening model coupled with vibration parameters can provide valuable insights into the mechanical behavior of materials during vibration-assisted deformation.
Due to the magnitude of *S* is determined by the amplitude *A*. By varying the amplitude, we can observe the corresponding changes in the vibration hardening value. Therefore, it can be deduced that *f*(*A*) also depends on the amplitude *A*. The specific expression for *f*(*A*) is as follows:

$$f(A) = \left\{ k \times \left( \frac{A}{A_0} \right)^q + 1 \right\}^m \qquad (6)$$

where *A* denotes the amplitude; *A*₀ refers to reference amplitude (whose value is 1mm); *k* is constants; *m* and *q* are constants. Obviously, when *A* = 0, *f* (0) = 1. The σmax/σwv - A/A₀ relationship was established and shown in Fig. 5. According to the fitting results, a mathematical expression about *f* (*A*) was obtained:

$$f(A) = (2.27 \times A^{0.66} + 1)^{0.11} \qquad (7)$$

<figure><img src="image_670.png" /><figcaption>Fig. 5 Linear fitting results of σmax/σwv-A/A₀</figcaption></figure>

The fitting results of the modified Hockett-Sherby hardening model parameters are summarized in Table 4.

<table><caption>Table 4. Parameters of modified Hockett-Sherby model</caption><thead><tr><th>parameters</th><th>σ<sub>s</sub>(MPa)</th><th>σ<sub>i</sub>(MPa)</th><th>a</th><th>n</th><th>k</th><th>q</th></tr></thead><tbody><tr><td>value</td><td>336.14</td><td>212.44</td><td>12.09</td><td>0.87</td><td>2.27</td><td>0.66</td></tr><tr><th>parameters</th><th>m</th><th>c</th><th>E(GPa)</th><th>L(mm)</th><th>A<sub>0</sub>(mm)</th><th></th></tr><tr><td>value</td><td>0.107</td><td>1.09</td><td>45</td><td>50</td><td>1</td><td></td></tr></tbody></table>

In summary, when low-frequency axial vibration was applied to AZ31B magnesium alloy, the curve of stress variation with strain of the material can be described by equation (8):

$$ \begin{cases} \sigma_{max} = \{336.14 - 123.7 \times \exp(-12.09 \times \varepsilon^{0.87})\} \times (2.27 \times A^{0.66} + 1)^{0.11} \\ \sigma_{min} = 336.14 - 123.7 \times \exp(-12.09 \times \varepsilon^{0.87}) \times (2.27 \times A^{0.66} + 1)^{0.11} - \Delta\sigma \\ \Delta\sigma = 981 \times A \end{cases} \qquad (8) $$

To verify the accuracy of the model, the LFVT results with an amplitude of 0.06 mm and a frequency of 20 Hz were compared with the predicted results. The comparison result is shown in Fig. 6. The results demonstrated that the model accurately captures the mechanical behavior of AZ31B under LFVT. This suggests that the Hockett-Sherby hardening model coupled with vibration parameters can provide valuable insights into the mechanical behavior of materials during vibration-assisted deformation.
<figure><img src="image_660.png" /><figcaption>Fig. 6 Comparison between predicted results of models and experimental results.</figcaption></figure>

## Conclusion

This study conducted LFVT tests on AZ31B sheet and investigated the effects of amplitude and frequency on mechanical behavior. On this basis, a Hockett-Sherby model coupled with vibration parameters was established and validated. The main conclusions are as follows:

(1) With superimposed low-frequency vibration field, the flow stress curve has the characteristic of stress oscillation phenomenon caused by periodic loading and unloading. The value of stress fluctuation increases with the increasing amplitude. For AZ31B magnesium alloy, low-frequency vibration field causes vibration hardening, and the variation trend is basically consistent with the fluctuation stress value. Compared with amplitude, vibration frequency has less effect on the mechanical behavior.

(2) Within the framework of Hockett-Sherby hardening model, the impact of amplitude was involved, and the parameters were determined through statistical analysis of the aligning quasi-static tensile stress-strain curve, stress fluctuation value, and vibration hardening value. By taking the amplitude effect into account, the Hockett-Sherby hardening model effectively forecasts the mechanical behavior of AZ31B under LFVT.

## References

[1] T. Maeno, K. Osakada. K. Mori, Reduction of friction in compression of plates by load pulsation[J]. Int. J. Mach. Tool. Manu., 51 (2011) 612-617.
https://doi.org/10.1016/j.ijmachtools.2011.03.007

[2] Z. Yao, G. Kim, L. Faidley, L. Faidley, Q. Zou, D. Mei, Z. Chen, Effects of superimposed high-frequency vibration on deformation of aluminum in micro/meso-scale upsetting[J]. J. Mater. Process. Tech., 212 (2012) 640-646. https://doi.org/10.1016/j.jmatprotec.2011.10.017

[3] M. Cao, H. Hu, X. Jia, S. Tian, C. Zhao, X. Han, Mechanism of ultrasonic vibration assisted upsetting of 6061 aluminum alloy[J]. J. Manuf. Process., 59 (2020) 690-697.
https://doi.org/10.1016/j.jmapro.2020.09.070

[4] G. Han, W. Wan, Z. Zhang, L. Xu, F. Liu, H. Zhang, Experimental investigation into effects of different ultrasonic vibration modes in micro-extrusion process[J]. J. Manuf. Process., 67 (2021) 427-437. https://doi.org/10.1016/j.jmapro.2021.05.007

[5] L. Deng, P. Li, X. Wang, M. Zhang, J. Li, Influence of low-frequency vibrations on the compression behavior and microstructure of T2 copper[J]. Mater. Sci. Eng. A, 710 (2018) 129-135. https://doi.org/10.1016/j.msea.2017.10.083

[6] J. Lin, C. Pruncu, L. Zhu, J. Li, Y. Zhai, L. Chen, Y. Guan, G. Zhao, Deformation behavior and microstructure in the low-frequency vibration upsetting of titanium alloy[J]. J. Mater. Process. Tech., 299 (2022) 117360. https://doi.org/10.1016/j.jmatprotec.2021.117360
<figure><img src="image_671.png" /><figcaption>Fig. 6 Comparison between predicted results of models and experimental results.</figcaption></figure>

## Conclusion

This study conducted LFVT tests on AZ31B sheet and investigated the effects of amplitude and frequency on mechanical behavior. On this basis, a Hockett-Sherby model coupled with vibration parameters was established and validated. The main conclusions are as follows:

(1) With superimposed low-frequency vibration field, the flow stress curve has the characteristic of stress oscillation phenomenon caused by periodic loading and unloading. The value of stress fluctuation increases with the increasing amplitude. For AZ31B magnesium alloy, low-frequency vibration field causes vibration hardening, and the variation trend is basically consistent with the fluctuation stress value. Compared with amplitude, vibration frequency has less effect on the mechanical behavior.

(2) Within the framework of Hockett-Sherby hardening model, the impact of amplitude was involved, and the parameters were determined through statistical analysis of the aligning quasi-static tensile stress-strain curve, stress fluctuation value, and vibration hardening value. By taking the amplitude effect into account, the Hockett-Sherby hardening model effectively forecasts the mechanical behavior of AZ31B under LFVT.

## References

[1] T. Maeno, K. Osakada. K. Mori, Reduction of friction in compression of plates by load pulsation[J]. Int. J. Mach. Tool. Manu., 51 (2011) 612-617.
https://doi.org/10.1016/j.ijmachtools.2011.03.007

[2] Z. Yao, G. Kim, L. Faidley, L. Faidley, Q. Zou, D. Mei, Z. Chen, Effects of superimposed high-frequency vibration on deformation of aluminum in micro/meso-scale upsetting[J]. J. Mater. Process. Tech., 212 (2012) 640-646. https://doi.org/10.1016/j.jmatprotec.2011.10.017

[3] M. Cao, H. Hu, X. Jia, S. Tian, C. Zhao, X. Han, Mechanism of ultrasonic vibration assisted upsetting of 6061 aluminum alloy[J]. J. Manuf. Process., 59 (2020) 690-697.
https://doi.org/10.1016/j.jmapro.2020.09.070

[4] G. Han, W. Wan, Z. Zhang, L. Xu, F. Liu, H. Zhang, Experimental investigation into effects of different ultrasonic vibration modes in micro-extrusion process[J]. J. Manuf. Process., 67 (2021) 427-437. https://doi.org/10.1016/j.jmapro.2021.05.007

[5] L. Deng, P. Li, X. Wang, M. Zhang, J. Li, Influence of low-frequency vibrations on the compression behavior and microstructure of T2 copper[J]. Mater. Sci. Eng. A, 710 (2018) 129-135. https://doi.org/10.1016/j.msea.2017.10.083

[6] J. Lin, C. Pruncu, L. Zhu, J. Li, Y. Zhai, L. Chen, Y. Guan, G. Zhao, Deformation behavior and microstructure in the low-frequency vibration upsetting of titanium alloy[J]. J. Mater. Process. Tech., 299 (2022) 117360. https://doi.org/10.1016/j.jmatprotec.2021.117360
[7] B. Meng, B. Cao, M. Wan, C. Wang, D. Shan, Constitutive behavior and microstructural evolution in ultrasonic vibration assisted deformation of ultrathin superalloy sheet[J]. Int. J. Mech. Sci., 157 (2019) 609-618. https://doi.org/10.1016/j.ijmecsci.2019.05.009

[8] Z. Yao, G. Kim, Z. Wang, L. Faidley, Q. Zou, D. Mei, Z. Chen, Acoustic softening and residual hardening in aluminum: modeling and experiments[J]. Int J Plast., 39 (2012) 75-87. https://doi.org/10.1016/j.ijplas.2012.06.003

[9] Z. Xie, Y. Guan, J. Lin, J. Zhai, L. Zhu, Constitutive model of 6063 aluminum alloy under the ultrasonic vibration upsetting based on Johnson-Cook model[J]. Ultrasonics, 96 (2019) 1-9. https://doi.org/10.1016/j.ultras.2019.03.017

[10] A. Prabhakar, G.Verma, K. Hariharan, P. Pandey, M. Lee, S. Suwas, Dislocation density based constitutive model for ultrasonic assisted deformation[J]. Mech. Res. Commun. 85 (2017) 76-80. https://doi.org/10.1016/j.mechrescom.2017.08.003

[11] W. Zhang, Y. Xu, Q. Li, X. Zhuang, Z. Zhao, Constitutive modeling and deformation analysis of W-temper and peak aged 7075 alloy sheets under low frequency vibration assisted tension[J]. Int. J. Mech. Mater. Des., 19 (2023) 583-604. https://doi.org/10.1007/s10999-023-09647-8

[12] Y. Xu, X. Zhuang, W. Zhang, Q. Li, Z. Zhao, Mechanical behaviors and microstructure characteristics of W-tempered and peak-aged 7075 alloy sheets under low frequency vibration-assisted tension[J]. Mater. Sci. Eng. A, 833 (2022) 142338. https://doi.org/10.1016/j.msea.2021.142338

[13] G. Nevill, F. Brotzen, The effect of vibrations on the static yield strength of a low-carbon steel[J]. Proceedings of American Society for Testing Material, 57 (1957) 751-758.

[14] P. Li, X. Wang, M. Zhang, L. Deng, J. Jin, Compression deformation behavior and size effect of copper under low-frequency vibration [J]. Forging & Stamping Technology. 42 (2017) 140-145

[15] D. Meng, J. Ma, X. Zhao, Y. Guo, C. Zhu, M. Yu, Mechanical behavior and material property of low-carbon steel undergoing low-frequency vibration-assisted upsetting[J]. J. Mater. Res. Technol., 16(2022) 1846-1855. https://doi.org/10.1016/j.jmrt.2021.12.113

[16] H. Kirchner, W. Kromp, F. Prinz, P. Trimmel, Plastic deformation under simultaneous cyclic and unidirectional loading at low and ultrasonic frequencies[J]. Mater. Sci. Eng., 68 (1985) 197-206. https://doi.org/10.1016/0025-5416(85)90409-4
[7] B. Meng, B. Cao, M. Wan, C. Wang, D. Shan, Constitutive behavior and microstructural evolution in ultrasonic vibration assisted deformation of ultrathin superalloy sheet[J]. Int. J. Mech. Sci., 157 (2019) 609-618. https://doi.org/10.1016/j.ijmecsci.2019.05.009

[8] Z. Yao, G. Kim, Z. Wang, L. Faidley, Q. Zou, D. Mei, Z. Chen, Acoustic softening and residual hardening in aluminum: modeling and experiments[J]. Int J Plast., 39 (2012) 75-87. https://doi.org/10.1016/j.ijplas.2012.06.003

[9] Z. Xie, Y. Guan, J. Lin, J. Zhai, L. Zhu, Constitutive model of 6063 aluminum alloy under the ultrasonic vibration upsetting based on Johnson-Cook model[J]. Ultrasonics, 96 (2019) 1-9. https://doi.org/10.1016/j.ultras.2019.03.017

[10] A. Prabhakar, G. Verma, K. Hariharan, P. Pandey, M. Lee, S. Suwas, Dislocation density based constitutive model for ultrasonic assisted deformation[J]. Mech. Res. Commun. 85 (2017) 76-80. https://doi.org/10.1016/j.mechrescom.2017.08.003

[11] W. Zhang, Y. Xu, Q. Li, X. Zhuang, Z. Zhao, Constitutive modeling and deformation analysis of W-temper and peak aged 7075 alloy sheets under low frequency vibration assisted tension[J]. Int. J. Mech. Mater. Des., 19 (2023) 583-604. https://doi.org/10.1007/s10999-023-09647-8

[12] Y. Xu, X. Zhuang, W. Zhang, Q. Li, Z. Zhao, Mechanical behaviors and microstructure characteristics of W-tempered and peak-aged 7075 alloy sheets under low frequency vibration-assisted tension[J]. Mater. Sci. Eng. A, 833 (2022) 142338. https://doi.org/10.1016/j.msea.2021.142338

[13] G. Nevill, F. Brotzen, The effect of vibrations on the static yield strength of a low-carbon steel[J]. Proceedings of American Society for Testing Material, 57 (1957) 751-758.

[14] P. Li, X. Wang, M. Zhang, L. Deng, J. Jin, Compression deformation behavior and size effect of copper under low-frequency vibration [J]. Forging & Stamping Technology. 42 (2017) 140-145

[15] D. Meng, J. Ma, X. Zhao, Y. Guo, C. Zhu, M. Yu, Mechanical behavior and material property of low-carbon steel undergoing low-frequency vibration-assisted upsetting[J]. J. Mater. Res. Technol., 16(2022) 1846-1855. https://doi.org/10.1016/j.jmrt.2021.12.113

[16] H. Kirchner, W. Kromp, F. Prinz, P. Trimmel, Plastic deformation under simultaneous cyclic and unidirectional loading at low and ultrasonic frequencies[J]. Mater. Sci. Eng., 68 (1985) 197-206. https://doi.org/10.1016/0025-5416(85)90409-4
# Effect of ball milling processing on mechanical properties of extruded aluminum-graphene-composites with commercial and self-synthesized graphene sources

NEGENDANK Maik<fnref n="1" />,ᵃ*, HOERNSCHEMEYER Jens<fnref n="1" />,ᵇ, JAIN Nishant<fnref n="2" />,ᶜ and MUELLER Soeren<fnref n="1" />,ᵈ

**Keywords:** Aluminum Graphene Composites, Oxalic Acid Exfoliated Graphite, Ball Milling, Extrusion, Mechanical Properties

**Abstract.** In this study the effect of ball milling processing on the mechanical properties of extruded aluminum-graphene-composites was investigated. A commercial and a self-synthesized graphene source was applied respectively. It was found that rods of extruded high speed ball milled (HSBM) materials showed bad surface quality with massive cracks. The extrusion loads in indirect extrusion were 100% higher for HSBM material compared to material that was ball milled at lower rotation speed (LSBM). Investigations of mechanical properties revealed that for HSBM material TYS was increased 92% and UTS 118% compared to LSBM. Microhardness was also found to increase by up to 210% for HSBM material containing 1% graphene. However, since pure aluminum processed under same conditions also featured a drastic increase in hardness of 167%, it can be concluded that work hardening of the pure aluminum matrix seems to be the main strengthening mechanism. Furthermore, graphene agglomerates could be found locally in all extruded samples.

## Introduction

Graphene is a planar material consisting of sp<fnref n="2" />-hybridized carbon layers. Its first synthesis was reported in the year 2004 [1]. Due to the exceptional properties of monolayer graphene, such as extremely high strength (130GPa) and elastic modulus (1TPa), graphene is very attractive to be applied as strengthening additive for Metal-Matrix-Composites (MMCs) [2]. Problematic is the fact that due to strong Van der Waals forces, graphene layers tend to agglomerate [3]. However, significant improvements in mechanical properties of composites can only be achieved when graphene is well exfoliated, distributed homogeneously in the matrix and the graphene source is of good quality (low number of layers and defects). In a previous study the authors applied short duration disc milling as a time and thus cost-efficient approach to manufacture aluminum-graphene-composites [4]. The results showed increased TYS by 16%, UTS by 27% and microhardness of up to 49%. However, in the literature several studies on ball milling as a production step for such composites can be found. Often ball milling was found to be beneficial for dispersion of graphene nano particles when the milling parameters are chosen accordingly [5]. In ball milling the radial component of the collision force (the compression force) of the milling balls with the aluminum powder particles can result in their flattening, fragmenting or cold-welding. While the tangential component of the collision force (the shearing force) can contribute to a better dispersion of GNPs as graphene layers that are attached to another initially, might be separated through shearing [6]. Zhang et al. [7] applied ball milling with a rotation speed of 350rpm

<figure><img src="image_661.png" /></figure>

Content from this work may be used under the terms of the Creative Commons Attribution 3.0 license. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. Published under license by Materials Research Forum LLC.

<footnote n="1">¹Extrusion Research and Development Center, TU Berlin, Gustav-Meyer-Allee 25, 13355 Berlin, Germany</footnote>

<footnote n="2">²Chair of Advanced Ceramic Materials, Hardenbergstr. 40, 10623 Berlin, Germany</footnote>

<footnote>ᵃmaik.negendank@strangpressen.berlin.de, ᵇjens.hoernschemeyer@campus.tu-berlin.de,</footnote>

<footnote>ᶜnishant.jain@ceramics.tu-berlin.de, ᵈsoeren.mueller@tu-berlin.de</footnote>
# Effect of ball milling processing on mechanical properties of extruded aluminum-graphene-composites with commercial and self-synthesized graphene sources

NEGENDANK Maik<fnref n="1" />,ᵃ*, HOERNSCHEMEYER Jens<fnref n="1" />,ᵇ, JAIN Nishant<fnref n="2" />,ᶜ and MUELLER Soeren<fnref n="1" />,ᵈ

**Keywords:** Aluminum Graphene Composites, Oxalic Acid Exfoliated Graphite, Ball Milling, Extrusion, Mechanical Properties

**Abstract.** In this study the effect of ball milling processing on the mechanical properties of extruded aluminum-graphene-composites was investigated. A commercial and a self-synthesized graphene source was applied respectively. It was found that rods of extruded high speed ball milled (HSBM) materials showed bad surface quality with massive cracks. The extrusion loads in indirect extrusion were 100% higher for HSBM material compared to material that was ball milled at lower rotation speed (LSBM). Investigations of mechanical properties revealed that for HSBM material TYS was increased 92% and UTS 118% compared to LSBM. Microhardness was also found to increase by up to 210% for HSBM material containing 1% graphene. However, since pure aluminum processed under same conditions also featured a drastic increase in hardness of 167%, it can be concluded that work hardening of the pure aluminum matrix seems to be the main strengthening mechanism. Furthermore, graphene agglomerates could be found locally in all extruded samples.

## Introduction

Graphene is a planar material consisting of sp<fnref n="2" />-hybridized carbon layers. Its first synthesis was reported in the year 2004 [1]. Due to the exceptional properties of monolayer graphene, such as extremely high strength (130GPa) and elastic modulus (1TPa), graphene is very attractive to be applied as strengthening additive for Metal-Matrix-Composites (MMCs) [2]. Problematic is the fact that due to strong Van der Waals forces, graphene layers tend to agglomerate [3]. However, significant improvements in mechanical properties of composites can only be achieved when graphene is well exfoliated, distributed homogeneously in the matrix and the graphene source is of good quality (low number of layers and defects). In a previous study the authors applied short duration disc milling as a time and thus cost-efficient approach to manufacture aluminum-graphene-composites [4]. The results showed increased TYS by 16%, UTS by 27% and microhardness of up to 49%. However, in the literature several studies on ball milling as a production step for such composites can be found. Often ball milling was found to be beneficial for dispersion of graphene nano particles when the milling parameters are chosen accordingly [5]. In ball milling the radial component of the collision force (the compression force) of the milling balls with the aluminum powder particles can result in their flattening, fragmenting or cold-welding. While the tangential component of the collision force (the shearing force) can contribute to a better dispersion of GNPs as graphene layers that are attached to another initially, might be separated through shearing [6]. Zhang et al. [7] applied ball milling with a rotation speed of 350rpm

<figure><img src="image_672.png" /></figure>

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<footnote n="1">¹Extrusion Research and Development Center, TU Berlin, Gustav-Meyer-Allee 25, 13355 Berlin, Germany</footnote>

<footnote n="2">²Chair of Advanced Ceramic Materials, Hardenbergstr. 40, 10623 Berlin, Germany</footnote>

<footnote>ᵃmaik.negendank@strangpressen.berlin.de, ᵇjens.hoernschemeyer@campus.tu-berlin.de,</footnote>

<footnote>ᶜnishant.jain@ceramics.tu-berlin.de, ᵈsoeren.mueller@tu-berlin.de</footnote>
for up to 24h with a ball-to-powder weight ratio of 20:1. After extrusion at 300°C with an extrusion ratio of R=10:1 TYS as well as in UTS were improved by 50%. Jiang et al. [5] performed shift-speed ball milling consisting of a long-term milling step at low rotation speeds of 200rpm and a short-term step at 500rpm. Again, a ball-to-powder ratio of 20:1 was selected, and results showed improvements in TYS of 24% and in UTS of 20%. Zheng et al. [8] achieved 166% higher TYS and 116% higher UTS by applying 10h of low milling speeds (100rpm) followed by 2h at 300rpm (ball-to-powder ratio 16:1) and extrusion at 450°C with R=25:1.

In the current study aluminum-graphene-composites with commercial as well as self-synthesized graphene sources should be manufactured via powder metallurgical route. Mechanical stirring, ball milling and indirect extrusion processing were applied to investigate the effects of processing conditions on the composites' mechanical properties. Another aim was to study the differences in composite properties when commercially available GNPs or a graphene source synthesized by the authors was applied respectively.

## Materials and Methods

In this study gas atomized pure aluminum (99.7%) powder was applied as matrix material. It was provided by TLS Technik GmbH & Co. KG, Germany. Particle size distribution was given in a prior study [4], with mean particle size being D$_{50}$=49µm. Two different graphene sources were applied. The first was commercial graphene nano platelets (GNP) provided by Alfa Aesar® (AA-GNPs). The second source was graphene that was synthesized by the chair of advanced ceramic materials of TU Berlin (Prof. Gurlo). Synthesis applied disc milling of fine graphite with oxalic acid. Hence, throughout this study the second source will be named oxalic acid exfoliated graphite (OAEG). A detailed study on the synthesis and characterization of OAEG will be published elsewhere soon. AA-GNPs have already been characterized in [4].

Extrusion billets for the current study here were prepared via powder metallurgical processing route. In dry (as-received) condition graphene particles formed agglomerates due to strong Van der Waals forces. Hence, in a first step graphene was soluted in a flask with ethanol. To improve graphene exfoliation and reduce number and size of agglomerates, the solution was sonicated with a sonification finger for 10min during simultaneous electromagnetic stirring. 1.5wt.% stearic acid as a process control agent was applied and together with pure dry aluminum powder was filled into an Eirich mixer type EL1. Afterwards, the liquid graphene/ethanol mixture was slowly added. Batches with graphene contents of 0wt.%, 0.25wt.%, 0.5wt.% and 1.0wt.% were produced with the commercial AA-GNPs. Regarding the OAEG, contents of up to 0.5wt.% were prepared only, due to insufficient amount of OAEG-powder. Additionally, prior results indicated decreasing properties for graphene contents above 0.5wt.% [4, 9]. Then the powder mixture was stirred for 6min at a rotation speed of 1800rpm. The mixed powders were then dried in a vacuum furnace at 50°C for 8h. After that powders were ball milled in a planetary mill (Fritsch Pulverisette 5) at different milling durations (1h, 4h, 8h) and rotation speeds (100rpm, 300rpm). Steel balls with a diameter of 12mm were applied for milling at a ball to powder weight ratio of 10:1. At the high rotation speed of 300rpm the milling process had to be conducted in stints due to strong heat generation. Hence, milling was conducted for 15min followed by a pause of 15min respectively. After ball milling, powders were degassed in a chamber furnace at 400°C in argon atmosphere to remove the stearic acid. Powder particle morphology in dependence of ball milling conditions were investigated in SEM (JEOL 640). Particle size could not yet be determined due to a technical defect of the particle size analyzer.

Milled powders were compacted into discs in a uniaxial press at room temperature. 32g powder were used for one each disc, resulting in a height of approximately 20mm. Compaction was conducted with a force of 150kN (pressure of 230MPa).

Extrusion processing was conducted in indirect extrusion mode on a 0.5MN extrusion press at the Extrusion Research and Development Center of TU Berlin. Three compacted discs were
for up to 24h with a ball-to-powder weight ratio of 20:1. After extrusion at 300°C with an extrusion ratio of R=10:1 TYS as well as in UTS were improved by 50%. Jiang et al. [5] performed shift-speed ball milling consisting of a long-term milling step at low rotation speeds of 200rpm and a short-term step at 500rpm. Again, a ball-to-powder ratio of 20:1 was selected, and results showed improvements in TYS of 24% and in UTS of 20%. Zheng et al. [8] achieved 166% higher TYS and 116% higher UTS by applying 10h of low milling speeds (100rpm) followed by 2h at 300rpm (ball-to-powder ratio 16:1) and extrusion at 450°C with R=25:1.

In the current study aluminum-graphene-composites with commercial as well as self-synthesized graphene sources should be manufactured via powder metallurgical route. Mechanical stirring, ball milling and indirect extrusion processing were applied to investigate the effects of processing conditions on the composites' mechanical properties. Another aim was to study the differences in composite properties when commercially available GNPs or a graphene source synthesized by the authors was applied respectively.

## Materials and Methods

In this study gas atomized pure aluminum (99.7%) powder was applied as matrix material. It was provided by TLS Technik GmbH & Co. KG, Germany. Particle size distribution was given in a prior study [4], with mean particle size being D$_{50}$=49µm. Two different graphene sources were applied. The first was commercial graphene nano platelets (GNP) provided by Alfa Aesar® (AA-GNPs). The second source was graphene that was synthesized by the chair of advanced ceramic materials of TU Berlin (Prof. Gurlo). Synthesis applied disc milling of fine graphite with oxalic acid. Hence, throughout this study the second source will be named oxalic acid exfoliated graphite (OAEG). A detailed study on the synthesis and characterization of OAEG will be published elsewhere soon. AA-GNPs have already been characterized in [4].

Extrusion billets for the current study here were prepared via powder metallurgical processing route. In dry (as-received) condition graphene particles formed agglomerates due to strong Van der Waals forces. Hence, in a first step graphene was soluted in a flask with ethanol. To improve graphene exfoliation and reduce number and size of agglomerates, the solution was sonicated with a sonification finger for 10min during simultaneous electromagnetic stirring. 1.5wt.% stearic acid as a process control agent was applied and together with pure dry aluminum powder was filled into an Eirich mixer type EL1. Afterwards, the liquid graphene/ethanol mixture was slowly added. Batches with graphene contents of 0wt.%, 0.25wt.%, 0.5wt.% and 1.0wt.% were produced with the commercial AA-GNPs. Regarding the OAEG, contents of up to 0.5wt.% were prepared only, due to insufficient amount of OAEG-powder. Additionally, prior results indicated decreasing properties for graphene contents above 0.5wt.% [4, 9]. Then the powder mixture was stirred for 6min at a rotation speed of 1800rpm. The mixed powders were then dried in a vacuum furnace at 50°C for 8h. After that powders were ball milled in a planetary mill (Fritsch Pulverisette 5) at different milling durations (1h, 4h, 8h) and rotation speeds (100rpm, 300rpm). Steel balls with a diameter of 12mm were applied for milling at a ball to powder weight ratio of 10:1. At the high rotation speed of 300rpm the milling process had to be conducted in stints due to strong heat generation. Hence, milling was conducted for 15min followed by a pause of 15min respectively. After ball milling, powders were degassed in a chamber furnace at 400°C in argon atmosphere to remove the stearic acid. Powder particle morphology in dependence of ball milling conditions were investigated in SEM (JEOL 640). Particle size could not yet be determined due to a technical defect of the particle size analyzer.

Milled powders were compacted into discs in a uniaxial press at room temperature. 32g powder were used for one each disc, resulting in a height of approximately 20mm. Compaction was conducted with a force of 150kN (pressure of 230MPa).

Extrusion processing was conducted in indirect extrusion mode on a 0.5MN extrusion press at the Extrusion Research and Development Center of TU Berlin. Three compacted discs were
stacked respectively for one billet. Discs and extrusion tools were heated to extrusion temperature of 350°C in a chamber furnace. Soaking time for Al-graphene-compacts was 45min to 60min. The container with diameter of 30mm was heated to the same temperature. A conic extrusion die with an orifice diameter of 8.0mm and a die angle of 2α=150° was applied for the extrusions. Ram speed was set to 3mm/s. For each condition (GNP-type, GNP-content, milling duration and milling speed) only one rod was extruded respectively.

After extrusion rod surface quality was evaluated and compared. Then, sections for mechanical testing were cut from the center of the rods towards the end. Tensile testing was performed on a MTS 810 testing machine at a strain rate of 3 · 10⁻⁴s⁻¹. Characterization of GNP-dispersion was conducted by optical microscopy on rod cross sections after etching with Dix-Keller's etchant.

## Results and discussion

The effects of ball milling parameters on aluminum as well as aluminum/graphene particles were investigated via SEM. The initial aluminum powder was of spherical shape [4]. Fig. 1 shows that after ball milling for 4h at 100rpm Al-particles were still predominantly round-shaped. However, their surface was rough and additionally some more plate-like shaped particles were also found. After a longer milling duration of 8h at 100rpm nearly all particles were deformed into plates. HSBM at 300rpm for 1h resulted in a significant reduction of the thickness of the plate-shaped particles as they were flattened significantly. After 4h of ball milling at 300rpm the particle morphology seemed to be unchanged but flattened aluminum particles seemed to be stacked and attached to each other. This indicates that cold-welding might have occurred.

<figure><img src="image_662.png" /><figcaption>Fig. 1: SEM images of particle size and morphology of pure aluminum powders in dependence of ball milling conditions</figcaption></figure>

Ball milled composite powders with 0.5% AA-GNPs and 0.5% OAEG were also investigated with SEM (Fig. 2). Graphene particles are indicated with arrows exemplarily. It was revealed that after ball milling AA-GNPs as well as OAEG were found to be attached to the surface of Al-particles. The commercial AA-GNPs seemed to be of bigger lateral size. Furthermore, they were found to be stacked on top of each other more often than the OAEG after same milling conditions. Hence, it can be concluded that OAEG are finer and better dispersed than AA-GNPs. Both factors should be favorable for improving mechanical properties of aluminum/graphene composites [3]. However, after 4h of HSBM at 300rpm, no graphene particles were found on the surface of aluminum powder particles. Hence, they must have been either reduced in size dramatically or dispersed into the aluminum particles during HSBM.
stacked respectively for one billet. Discs and extrusion tools were heated to extrusion temperature of 350°C in a chamber furnace. Soaking time for Al-graphene-compacts was 45min to 60min. The container with diameter of 30mm was heated to the same temperature. A conic extrusion die with an orifice diameter of 8.0mm and a die angle of $2\alpha=150°$ was applied for the extrusions. Ram speed was set to 3mm/s. For each condition (GNP-type, GNP-content, milling duration and milling speed) only one rod was extruded respectively.

After extrusion rod surface quality was evaluated and compared. Then, sections for mechanical testing were cut from the center of the rods towards the end. Tensile testing was performed on a MTS 810 testing machine at a strain rate of $3 \cdot 10^{-4} s^{-1}$. Characterization of GNP-dispersion was conducted by optical microscopy on rod cross sections after etching with Dix-Keller's etchant.

## Results and discussion

The effects of ball milling parameters on aluminum as well as aluminum/graphene particles were investigated via SEM. The initial aluminum powder was of spherical shape [4]. Fig. 1 shows that after ball milling for 4h at 100rpm Al-particles were still predominantly round-shaped. However, their surface was rough and additionally some more plate-like shaped particles were also found. After a longer milling duration of 8h at 100rpm nearly all particles were deformed into plates. HSBM at 300rpm for 1h resulted in a significant reduction of the thickness of the plate-shaped particles as they were flattened significantly. After 4h of ball milling at 300rpm the particle morphology seemed to be unchanged but flattened aluminum particles seemed to be stacked and attached to each other. This indicates that cold-welding might have occurred.

<figure><img src="image_673.png" /><figcaption>Fig. 1: SEM images of particle size and morphology of pure aluminum powders in dependence of ball milling conditions</figcaption></figure>

Ball milled composite powders with 0.5% AA-GNPs and 0.5% OAEG were also investigated with SEM (Fig. 2). Graphene particles are indicated with arrows exemplarily. It was revealed that after ball milling AA-GNPs as well as OAEG were found to be attached to the surface of Al-particles. The commercial AA-GNPs seemed to be of bigger lateral size. Furthermore, they were found to be stacked on top of each other more often than the OAEG after same milling conditions. Hence, it can be concluded that OAEG are finer and better dispersed than AA-GNPs. Both factors should be favorable for improving mechanical properties of aluminum/graphene composites [3]. However, after 4h of HSBM at 300rpm, no graphene particles were found on the surface of aluminum powder particles. Hence, they must have been either reduced in size dramatically or dispersed into the aluminum particles during HSBM.
<figure><img src="image_663.png" /><figcaption>Fig. 2: SEM pictures of 0.5% Al/graphene composite powders after ball milling with different milling parameters and degassing (graphene exemplarily marked with arrows)</figcaption></figure>

Next the effects of ball milling conditions (energy input), GNP-content and GNP-type on the extrusion load will be described. The extrusion loads (peak force at the beginning of the indirect extrusion process) are displayed in Fig. 3. It can be noted that the predominant factor for the extrusion load is the rotation speed during ball milling. At low rotation speeds of 100rpm extrusion forces were around 200kN. In fact, at the lower milling duration of 4h forces between 172kN (0.5% OAEG) and 191kN (0.25% OAEG) were observed. At longer milling duration of 8h, similar but slightly higher extrusion loads between 172kN (1.0% AA-GNPs) and 213kN (0.5% AA-GNPs) were found. The slightly increasing tendency could be the result of stronger work hardening of aluminum particles due to longer milling duration and thus increased material deformation. The effect of graphene content on extrusion load was not significant in these cases. On the other hand, for extrusion of the HSBM-materials a significant increase in extrusion load was found for all extrusions. For the 1h HSBM milled materials extrusion loads between 426kN (0.25% AA-GNPs) and 483kN (0.5kN) were found. Hence, extrusion loads doubled due to HSBM. For the longer milling duration of 4h at 300rpm the results tended to decrease. Extrusion loads between 309kN (0.5% AA-GNP) and 388kN (pure Al) were measured in this case.

<figure><img src="image_664.png" /><figcaption>Fig. 3: Effect of ball milling conditions, graphene type and graphene content on peak extrusion loads in indirect extrusion processing</figcaption></figure>

After extrusion the quality of extruded rods was inspected. Rods of low-speed ball milled material (LSBM) (milled at 100rpm) showed some grooves only. However, all rods of the HSBM-
<figure><img src="image_674.png" /><figcaption>Fig. 2: SEM pictures of 0.5% Al/graphene composite powders after ball milling with different milling parameters and degassing (graphene exemplarily marked with arrows)</figcaption></figure>

Next the effects of ball milling conditions (energy input), GNP-content and GNP-type on the extrusion load will be described. The extrusion loads (peak force at the beginning of the indirect extrusion process) are displayed in Fig. 3. It can be noted that the predominant factor for the extrusion load is the rotation speed during ball milling. At low rotation speeds of 100rpm extrusion forces were around 200kN. In fact, at the lower milling duration of 4h forces between 172kN (0.5% OAEG) and 191kN (0.25% OAEG) were observed. At longer milling duration of 8h, similar but slightly higher extrusion loads between 172kN (1.0% AA-GNPs) and 213kN (0.5% AA-GNPs) were found. The slightly increasing tendency could be the result of stronger work hardening of aluminum particles due to longer milling duration and thus increased material deformation. The effect of graphene content on extrusion load was not significant in these cases. On the other hand, for extrusion of the HSBM-materials a significant increase in extrusion load was found for all extrusions. For the 1h HSBM milled materials extrusion loads between 426kN (0.25% AA-GNPs) and 483kN (0.5kN) were found. Hence, extrusion loads doubled due to HSBM. For the longer milling duration of 4h at 300rpm the results tended to decrease. Extrusion loads between 309kN (0.5% AA-GNP) and 388kN (pure Al) were measured in this case.

<figure><img src="image_675.png" /><figcaption>Fig. 3: Effect of ball milling conditions, graphene type and graphene content on peak extrusion loads in indirect extrusion processing</figcaption></figure>

After extrusion the quality of extruded rods was inspected. Rods of low-speed ball milled material (LSBM) (milled at 100rpm) showed some grooves only. However, all rods of the HSBM-
material including those of pure aluminum featured massive surface cracks. Fig. 4a visualizes the effect of ball milling conditions for rods with 0.5% AA-GNP content. Fig. 4b on the other hand illustrates the effect of different AA-GNP contents on rod surface quality for composite materials that were ball milled for 8h at 100rpm. It can be noted that rod surface quality decreased with increasing GNP content. This effect was also observed in previous studies for disc milled and indirectly extruded Al/GNP composites [4] as well as for mechanically mixed and directly extruded composites [9].

<figure><img src="image_665.png" /><figcaption>Fig. 4: Rod surface quality in dependence of a) applied ball milling conditions for 0.5% AA-GNP composites b) AA-GNP content (ball milled for 8h at 100rpm)</figcaption></figure>

To characterize the effect of different ball milling parameters on the composites' mechanical properties, tensile tests as well as microhardness testing were performed. As noted previously, due to surface cracking, samples for tensile testing could mostly be prepared for LSBM-material only. The tensile yield strengths (TYS) as well as the ultimate tensile strengths (UTS) and fracture strains (E<sub>fr</sub>) are summarized in Table 2 and visualized in Figures 5a and 5b respectively.

<table><caption>Table 2: Summary of tensile test results</caption><thead><tr><th rowspan="2">ball milling conditions</th><th colspan="3">pure Al</th><th colspan="3">0.25% AA-GNPs</th><th colspan="3">0.5% AA-GNPs</th><th colspan="3">1.0% AA-GNPs</th><th colspan="3">0.25% OAEG</th><th colspan="3">0.5% OAEG</th></tr><tr><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th></tr><tr><th></th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th></tr></thead><tbody><tr><td>4h, 100rpm</td><td>81±1</td><td>114±1</td><td>36±6</td><td>78±1</td><td>112±1</td><td>30±5</td><td>82±1</td><td>114±1</td><td>24±1</td><td>79±4</td><td>112±1</td><td>22±3</td><td>85±2</td><td>118±2</td><td>21±1</td><td>80±1</td><td>115±1</td><td>28±1</td><td></td><td></td><td></td></tr><tr><td>8h, 100rpm</td><td>90±6</td><td>125±1</td><td>28±1</td><td>84±1</td><td>124±1</td><td>24±1</td><td>86±1</td><td>128±1</td><td>25±3</td><td>88±4</td><td>124±1</td><td>20±1</td><td>88±1</td><td>127±1</td><td>25±3</td><td>86±1</td><td>125±1</td><td>27±1</td><td></td><td></td><td></td></tr><tr><td>1h, 300rpm</td><td>-</td><td>-</td><td>-</td><td>150±5</td><td>244±1</td><td>7±1</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td></tr></tbody></table>

For composites being ball milled at 100rpm for 4h prior to extrusion, TYS values between 78MPa and 85MPa were found. The pure aluminum reference featured a TYS of 81MPa. UTS values for these composites were between 112MPa (1% AA-GNP) and 118MPa (0.25% OAEG). The UTS of the pure Al sample was 114MPa.

Samples that were milled at 100rpm for a longer duration of 8h featured slightly higher values. The TYS values were found in the range between 84MPa (0.25% AA-GNPs) and 88MPa (0.25% OAEG) with pure Al at 90MPa. UTS values were between 124MPa and 128MPa with pure Al being at 125MPa. Thus, TYS increased by up to +11% and UTS by up to +7%. Further, it can be learned that for LSBM composites, the effect of graphene content as well as the applied graphene source (commercial AA-GNPs and self-synthesized OAEG) did not have a strong effect on TYS and UTS. Additionally, the strengths values of LSBM-material are similar to those gained after short term disc milling in a prior study [4].
material including those of pure aluminum featured massive surface cracks. Fig. 4a visualizes the effect of ball milling conditions for rods with 0.5% AA-GNP content. Fig. 4b on the other hand illustrates the effect of different AA-GNP contents on rod surface quality for composite materials that were ball milled for 8h at 100rpm. It can be noted that rod surface quality decreased with increasing GNP content. This effect was also observed in previous studies for disc milled and indirectly extruded Al/GNP composites [4] as well as for mechanically mixed and directly extruded composites [9].

<figure><img src="image_676.png" /><figcaption>Fig. 4: Rod surface quality in dependence of a) applied ball milling conditions for 0.5% AA-GNP composites b) AA-GNP content (ball milled for 8h at 100rpm)</figcaption></figure>

To characterize the effect of different ball milling parameters on the composites' mechanical properties, tensile tests as well as microhardness testing were performed. As noted previously, due to surface cracking, samples for tensile testing could mostly be prepared for LSBM-material only. The tensile yield strengths (TYS) as well as the ultimate tensile strengths (UTS) and fracture strains (E<sub>fr</sub>) are summarized in Table 2 and visualized in Figures 5a and 5b respectively.

<table><caption>Table 2: Summary of tensile test results</caption><thead><tr><th rowspan="2">ball milling conditions</th><th colspan="3">pure Al</th><th colspan="3">0.25% AA-GNPs</th><th colspan="3">0.5% AA-GNPs</th><th colspan="3">1.0% AA-GNPs</th><th colspan="3">0.25% OAEG</th><th colspan="3">0.5% OAEG</th></tr><tr><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th><th>TYS</th><th>UTS</th><th>E<sub>fr</sub></th></tr><tr><th></th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th><th colspan="3">[MPa]</th></tr></thead><tbody><tr><td>4h, 100rpm</td><td>81±1</td><td>114±1</td><td>36±6</td><td>78±1</td><td>112±1</td><td>30±5</td><td>82±1</td><td>114±1</td><td>24±1</td><td>79±4</td><td>112±1</td><td>22±3</td><td>85±2</td><td>118±2</td><td>21±1</td><td>80±1</td><td>115±1</td><td>28±1</td><td></td><td></td><td></td></tr><tr><td>8h, 100rpm</td><td>90±6</td><td>125±1</td><td>28±1</td><td>84±1</td><td>124±1</td><td>24±1</td><td>86±1</td><td>128±1</td><td>25±3</td><td>88±4</td><td>124±1</td><td>20±1</td><td>88±1</td><td>127±1</td><td>25±3</td><td>86±1</td><td>125±1</td><td>27±1</td><td></td><td></td><td></td></tr><tr><td>1h, 300rpm</td><td>-</td><td>-</td><td>-</td><td>150±5</td><td>244±1</td><td>7±1</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td></tr></tbody></table>

For composites being ball milled at 100rpm for 4h prior to extrusion, TYS values between 78MPa and 85MPa were found. The pure aluminum reference featured a TYS of 81MPa. UTS values for these composites were between 112MPa (1% AA-GNP) and 118MPa (0.25% OAEG). The UTS of the pure Al sample was 114MPa.

Samples that were milled at 100rpm for a longer duration of 8h featured slightly higher values. The TYS values were found in the range between 84MPa (0.25% AA-GNPs) and 88MPa (0.25% OAEG) with pure Al at 90MPa. UTS values were between 124MPa and 128MPa with pure Al being at 125MPa. Thus, TYS increased by up to +11% and UTS by up to +7%. Further, it can be learned that for LSBM composites, the effect of graphene content as well as the applied graphene source (commercial AA-GNPs and self-synthesized OAEG) did not have a strong effect on TYS and UTS. Additionally, the strengths values of LSBM-material are similar to those gained after short term disc milling in a prior study [4].
However, from the extruded rod containing 0.25% AA-GNPs that was initially HSBM for 1h at 300rpm, despite some minor surface cracks, it was possible to manufacture samples for tensile testing. For these samples significantly higher strengths of TYS=150MPa and UTS=244MPa were observed. Thus, the HSBM samples have a 92% higher TYS and a 118% higher UTS when being compared to the pure Al samples ball milled at 100rpm for 4h.

Regarding the fracture strain it was observed that the increases in strengths came at the cost of lower fracture strain and thus ductility. Composites with 0.25% AA-GNPs showed a slight decrease from E<sub>fr</sub>=30% to E<sub>fr</sub>=25% when LSBM duration was varied from 4h to 8h. More drastically the fracture strain reduced to E<sub>fr</sub>=7% when HSBM at 300rpm was conducted for 1h.

Sections from all extruded rods including those with severely cracked surfaces were tested to determine the microhardness. The results are summarized in Table 3 and visualized in Fig. 6. Very similar to the observed tendencies of tensile testing results, microhardness values for rods of LSBM-material were between 36HV to 37HV when milling duration was 4h. Increasing milling duration to 8h resulted in slightly higher microhardness values between 39 and 42HV, whereas 42HV was found for the 0.5% AA-GNP composite rod. There was no significant effect of GNP-type and GNP-content on the microhardness of LSBM material. Again, these values are in good comparison with prior results of short-duration disc milled Al/graphene-composites [4]. However, when being compared to values of 4h LSBM, processing via 4h of HSBM increased hardness values substantially. Improvements by up to 159% (in case of 0.5% AA-GNPs) to values between 75HV and 90HV (1h, 300rpm) and even further by up to 210% (in case of 1.0% AA-GNPs) to values between 95HV to 115HV in case of 4h HSBM were found. HSBM also significantly increased microhardness of pure aluminum by 167%. Hence, it might be concluded that work hardening of aluminum particles during the milling process was the main strengthening mechanism and not the addition of graphene particles. An upcoming study will focus on improving rod surface quality during extrusion. Then it can be verified, if the relationships between ball milling parameters and microhardness found in the current study show the same tendency in tensile tests.

<figure><img src="image_666.png" /><figcaption>Fig. 5: Effects of ball milling conditions, graphene content and graphene type on a) Tensile yield strength (TYS) b) Ultimate tensile strength (UTS) of Al/graphene composites</figcaption></figure>
However, from the extruded rod containing 0.25% AA-GNPs that was initially HSBM for 1h at 300rpm, despite some minor surface cracks, it was possible to manufacture samples for tensile testing. For these samples significantly higher strengths of TYS=150MPa and UTS=244MPa were observed. Thus, the HSBM samples have a 92% higher TYS and a 118% higher UTS when being compared to the pure Al samples ball milled at 100rpm for 4h.

Regarding the fracture strain it was observed that the increases in strengths came at the cost of lower fracture strain and thus ductility. Composites with 0.25% AA-GNPs showed a slight decrease from E<sub>fr</sub>=30% to E<sub>fr</sub>=25% when LSBM duration was varied from 4h to 8h. More drastically the fracture strain reduced to E<sub>fr</sub>=7% when HSBM at 300rpm was conducted for 1h.

Sections from all extruded rods including those with severely cracked surfaces were tested to determine the microhardness. The results are summarized in Table 3 and visualized in Fig. 6. Very similar to the observed tendencies of tensile testing results, microhardness values for rods of LSBM-material were between 36HV to 37HV when milling duration was 4h. Increasing milling duration to 8h resulted in slightly higher microhardness values between 39 and 42HV, whereas 42HV was found for the 0.5% AA-GNP composite rod. There was no significant effect of GNP-type and GNP-content on the microhardness of LSBM material. Again, these values are in good comparison with prior results of short-duration disc milled Al/graphene-composites [4]. However, when being compared to values of 4h LSBM, processing via 4h of HSBM increased hardness values substantially. Improvements by up to 159% (in case of 0.5% AA-GNPs) to values between 75HV and 90HV (1h, 300rpm) and even further by up to 210% (in case of 1.0% AA-GNPs) to values between 95HV to 115HV in case of 4h HSBM were found. HSBM also significantly increased microhardness of pure aluminum by 167%. Hence, it might be concluded that work hardening of aluminum particles during the milling process was the main strengthening mechanism and not the addition of graphene particles. An upcoming study will focus on improving rod surface quality during extrusion. Then it can be verified, if the relationships between ball milling parameters and microhardness found in the current study show the same tendency in tensile tests.

<figure><img src="image_677.png" /><figcaption>Fig. 5: Effects of ball milling conditions, graphene content and graphene type on a) Tensile yield strength (TYS) b) Ultimate tensile strength (UTS) of Al/graphene composites</figcaption></figure>
<table><caption>Table 3: Microhardness (Vickers) of extruded Al/graphene composites in dependence of ball milling conditions, values in HV0.5</caption><thead><tr><th rowspan="2">composite</th><th colspan="4">ball milling conditions</th></tr><tr><th>4h, 100rpm</th><th>8h, 100rpm</th><th>1h, 300rpm</th><th>4h, 300rpm</th></tr></thead><tbody><tr><td>pure Al</td><td>36±1</td><td>39±1</td><td>84±2</td><td>96±2</td></tr><tr><td>0.25% AA-GNPs</td><td>38±1</td><td>39±1</td><td>75±1</td><td>95±3</td></tr><tr><td>0.5% AA-GNPs</td><td>37±1</td><td>42±3</td><td>90±5</td><td>96±2</td></tr><tr><td>1.0% AA-GNPs</td><td>37±2</td><td>41±1</td><td>77±2</td><td>115±2</td></tr><tr><td>0.25% OAEG</td><td>37±1</td><td>39±1</td><td>-</td><td>-</td></tr><tr><td>0.5% OAEG</td><td>36±1</td><td>39±1</td><td>-</td><td>-</td></tr></tbody></table>

<figure><img src="image_667.png" /><figcaption>Fig. 6: Effects of ball milling conditions, graphene content and graphene type on rod microhardness</figcaption></figure>

Since a homogeneous dispersion of graphene particles in the pure-Al matrix is important for significant improvements in strengths of composites, the dispersion on extruded rod cross sections was investigated by optical microscopy. The results are given at high microscope magnification of 1000x in Fig. 7 and for overview purposes at a lower magnification of 50x in Fig. 8.

At higher magnification (1000x) graphene particles (dark particles in the images) seemed to be relatively well dispersed in the pure-Al matrix even for high contents of 1.0% AA-GNPs. For the rods of LSBM materials (100rpm) there was no significant effect of milling duration on the graphene dispersion. Furthermore, the two different graphene sources of AA-GNPs and OAEG seemed to be dispersed similarly at same contents. However, after HSBM for 1h graphene particles seemed to have aligned close to each other in a meandric shape. After 4h of HSBM graphene particles were not visible in matrix anymore. It seems that longer duration for HSBM of 4h at 300rpm reduced the size of graphene particles significantly, so that they are longer visible in optical microscopy.

Fig. 8 shows that at lower magnification graphene agglomerates (black particles) can be found locally in almost all of the rod cross sections. Agglomerates are known to be detrimental for mechanical properties of composites as they might act as sites for crack initiation under tensile loading. It has to be concluded that the processing route of sonication, intensive mechanical stirring, ball milling and extrusion was not sufficient to fully dissolve preexisting (due to Van der Waals forces) graphene agglomerates. Hence, in an upcoming study the effect of additional shear deformation on the graphene dispersion and the mechanical properties will be investigated. During significant shearing, graphene layers could potentially be “sheared off” from one another. Thus, shearing would be beneficial for graphene dispersion as agglomerates may be dissolved.
<table><caption>Table 3: Microhardness (Vickers) of extruded Al/graphene composites in dependence of ball milling conditions, values in HV0.5</caption><thead><tr><th rowspan="2">composite</th><th colspan="4">ball milling conditions</th></tr><tr><th>4h, 100rpm</th><th>8h, 100rpm</th><th>1h, 300rpm</th><th>4h, 300rpm</th></tr></thead><tbody><tr><td>pure Al</td><td>36±1</td><td>39±1</td><td>84±2</td><td>96±2</td></tr><tr><td>0.25% AA-GNPs</td><td>38±1</td><td>39±1</td><td>75±1</td><td>95±3</td></tr><tr><td>0.5% AA-GNPs</td><td>37±1</td><td>42±3</td><td>90±5</td><td>96±2</td></tr><tr><td>1.0% AA-GNPs</td><td>37±2</td><td>41±1</td><td>77±2</td><td>115±2</td></tr><tr><td>0.25% OAEG</td><td>37±1</td><td>39±1</td><td>-</td><td>-</td></tr><tr><td>0.5% OAEG</td><td>36±1</td><td>39±1</td><td>-</td><td>-</td></tr></tbody></table>

<figure><img src="image_678.png" /><figcaption>Fig. 6: Effects of ball milling conditions, graphene content and graphene type on rod microhardness</figcaption></figure>

Since a homogeneous dispersion of graphene particles in the pure-Al matrix is important for significant improvements in strengths of composites, the dispersion on extruded rod cross sections was investigated by optical microscopy. The results are given at high microscope magnification of 1000x in Fig. 7 and for overview purposes at a lower magnification of 50x in Fig. 8.

At higher magnification (1000x) graphene particles (dark particles in the images) seemed to be relatively well dispersed in the pure-Al matrix even for high contents of 1.0% AA-GNPs. For the rods of LSBM materials (100rpm) there was no significant effect of milling duration on the graphene dispersion. Furthermore, the two different graphene sources of AA-GNPs and OAEG seemed to be dispersed similarly at same contents. However, after HSBM for 1h graphene particles seemed to have aligned close to each other in a meandric shape. After 4h of HSBM graphene particles were not visible in matrix anymore. It seems that longer duration for HSBM of 4h at 300rpm reduced the size of graphene particles significantly, so that they are longer visible in optical microscopy.

Fig. 8 shows that at lower magnification graphene agglomerates (black particles) can be found locally in almost all of the rod cross sections. Agglomerates are known to be detrimental for mechanical properties of composites as they might act as sites for crack initiation under tensile loading. It has to be concluded that the processing route of sonication, intensive mechanical stirring, ball milling and extrusion was not sufficient to fully dissolve preexisting (due to Van der Waals forces) graphene agglomerates. Hence, in an upcoming study the effect of additional shear deformation on the graphene dispersion and the mechanical properties will be investigated. During significant shearing, graphene layers could potentially be “sheared off” from one another. Thus, shearing would be beneficial for graphene dispersion as agglomerates may be dissolved.
<figure><img src="image_668.png" /><figcaption>Fig. 7: OM-images of composite rod cross sections at microscope magnification of 1000x</figcaption></figure>

In order to investigate if during ball milling or extrusion processing oxidation of aluminum powder or formation of Al₄C₃ phase might have occurred, XRD phase analysis (measured with Bruker D8 advanced diffractometer) was conducted exemplarily. XRD was conducted on the sample with the highest strength (0.25% AA-GNPs, 1h, 300rpm) as well as on the sample with highest overall microhardness (1.0% AA-GNPs, 4h, 300rpm). The diffraction patterns are given in Fig. 9. In case of the sample with highest microhardness besides aluminum no other peaks were detected. In case of the specimen with highest microhardness peaks of aluminum oxide (cubic Al₂O₃) were found. Hence, it seems that at least in this case during HSBM for 1h with significant heat development in atmospheric environment, oxidation of aluminum powder occurred to some extent. The Al₄C₃-phase did not appear in the patterns and thus, did not seem to have formed during processing.