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ISBN print: 978-3-0364-0990-0
ISBN eBook: 978-3-0364-1990-9
DOI book: 10.4028/b-1S6kQl
© 2026 Trans Tech Publications Ltd, All Rights Reserved

Innovative Joining
by Forming

Edited by
Prof. Gianluca Buffa

Innovative Joining by Forming

Special topic volume with invited peer-reviewed papers only

Edited by
Prof. Gianluca Buffa

Scientific.Net

Copyright © 2026 Trans Tech Publications Ltd, Switzerland

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Volume 1185 of
Materials Science Forum
ISSN print 0255-5476
ISSN web 1662-9752

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp -1-0

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Accepted:
2026-04-08

Online:
2026-04-15

Buffa, Gianluca

Preface

The original version of this paper is available on https://www.scientific.net/MSF.1185.-1.pdf

The main topics of the issue are: mechanical fastening operations (clinching, riveting and so on); friction welding; friction stir welding and friction stir spot welding; hybrid joining processes; innovative joining processes (also using adhesives) including forming or pre-forming of the materials to be joined; computational methods for joining simulation of processes; applications of inverse analysis to material characterization in joining conditions; part integrity (prediction of residual stresses, distortions ecc); innovative measurement methods and devices to detect relevant variables in joining by forming (temperatures, pressures, vibrations, wear ....)

Editor

The original version of this paper is available on https://www.scientific.net/MSF.1185.-1.pdf

Co-editors

The original version of this paper is available on https://www.scientific.net/MSF.1185.-1.pdf

Merklein, Marion
Meschut, Gerson

University of Palermo

Friedrich-Alexander-UniversitätErlangen-Nürnberg Paderborn University

Table of Contents

Preface

Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Zhao Yi Liu, Li Yin, Wei Yang, Jing Hui Gao, Yang Lv, Zi Hao Wu, Chen Xi Wang

Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Daniel Yeniss Unni, Matteo Strano

Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

Jessica Sarris, Michael Lechner

Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Marco Zambelli, Sara Bocchi, Marco Negozio, Gianluca Danilo D'Urso, Claudio Giardini, Laura Antonini

Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Nik Schwichtenberg, Markus Tesch, Martina Elisabeth Sigl, Moritz Goeldner, Michael F. Zaeh

Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Aron Ringel, Ahmet Zahid Demirezen, Julika Hoyer, David Bailly, Simon Kammerloher, Wolfram Volk, Junhe Lian

A Hybrid Joining Strategy for Al-Cu Bimetallic Sheets Produced by Friction Stir Welding

Ioannis Papantoniou, Dimitrios E. Manolakos

Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

Armin Piwek, Julius Peddinghaus, Johanna Uhe, Kai Brunotte, Bernd Arno Behrens

Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Riccardo Puleo, Salvatore Russo, Gianluca Buffa, Livan Fratini

Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Austin Clark, Philip Barnett, Elizabeth Mamros, Brian Young, Ihab Ragai

Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Amith Tom Mathew, Armin Piwek, Jakub Mayer, Norman Mohnfeld, Bernd Arno Behrens, Johanna Uhe

Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Mohamad El Mehtedi, Pasquale Buonadonna, Barbara Reggiani, Riccardo Pelaccia, Lorenzo Donati, Gabriela Loi, Mauro Carta

Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Lasse Malaske, Uceu Suhuddin, Benjamin Klusemann

Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Pia Katharina Kaimann, Mathias Bobbert, Gerson Meschut

A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

Deekshith Reddy Devulapally, Thomas Tröster

Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Sheng Jing Liu, Yan Hong Mu, Xiao Hua Li, Hai Long Lei

Keywords Index

Authors Index

The original version of this eBook is available on https://www.scientific.net/book/innovative-joining-by-forming/978-3-0364-1990-9

Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 1-7

doi: 10.4028/p-8caZwL

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2025-12-21

Revised:
2026-02-15

Accepted:
2026-02-24

Online:
2026-04-15

Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

Zhaoyi Liu 1,a, Li Yin 1,b*, Wei Yang 1,c, Jinghui Gao 2,d, Yang Lv 1,e, Zihao Wu 3,f and Chenxi Wang 3,g

1 Beijing Institute of Smart Energy, Beijing, China

2 State Key Laboratory of Electrical Insulation and Power Equipment, Xi'an Jiaotong University, Xi'an, China

3 Power Research Institute, State Grid Shaanxi Power Co., Xi'an, China

a  markliu3@163.com, b  yinlibj@163.com, c  19630100@163.com,

Keywords: factory joint, XLPE insulation, interfacial bonding, ultrasonic vibration.

Abstract

The interface between the body and recovery insulation in factory joints is susceptible to localized electromechanical weaknesses, potentially compromising the long-term reliability of highvoltage direct current (HVDC) submarine cable systems. This study introduces ultrasonic vibration into the insulation recovery process of these joints. The effects of ultrasonic treatment are evaluated through tensile tests, thermal elongation tests, DC breakdown tests, and space-charge characterization. Results demonstrate that ultrasonic treatment significantly enhances interfacial tensile properties, improves deformation stability under thermal loading, increases the DC breakdown strength of the insulation interface, and significantly suppresses negative space-charge accumulation. These improvements are attributed to the high-frequency acoustic pressure and localized temperature rise induced by ultrasonic vibration, which promote molecular-chain melting and rearrangement, thereby reducing microscopic defects. This study provides an innovative method for the insulation recovery forming process in factory joints to improve interfacial reliability.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

Against the backdrop of global energy transition, the demand for long-distance and large-scale integration of renewable energy is steadily increasing. High voltage direct current (HVDC) transmission technology, with its advantages of low loss and large capacity, especially when combined with cross-linked polyethylene (XLPE) insulated cables, has become a core solution for offshore wind power integration and regional interconnections [1]. Nevertheless, the reliability of HVDC cable systems faces significant challenges, with factory joints representing key weak points [2,3].

During the manufacturing of factory joints, a "secondary bonding interface" forms between the cable body insulation and the recovery insulation [4]. Located within the thermal press cross-linking transition zone, this interface is susceptible to microscopic defects, uneven crystallization, and space charge accumulation. These issues lead to electric field distortion and accelerated insulation aging [5]. Prior studies indicate that controlling interface roughness and degassing processes can improve crystallization behavior and suppress charge accumulation [6]. In advanced manufacturing, ultrasonic technology has been applied in fields such as polymer welding to enhance interface properties. By generating high-frequency vibration-induced thermal and mechanical effects, ultrasonic treatment strengthens interface bonding and eliminates defects [7-11].

[picture]

Fig. 1. Schematic diagram of the factory joint structure in high-voltage submarine cables.

Fig. 1. Schematic diagram of the factory joint structure in high-voltage submarine cables.

In this study, ultrasonic vibration was introduced during the insulation recovery molding process of factory joints to enhance the secondary bonding at the interface. First, an ultrasonic vibration enhanced bonding platform was established. Subsequently, through tensile tests, thermal elongation tests, DC breakdown tests, and space charge tests, the changes in mechanical and electrical properties of the insulated interface under ultrasonic treatment were investigated. Finally, the mechanism underlying ultrasonic vibration enhancement was elucidated.

Materials and Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

An ultrasonic vibration-enhanced bonding platform was developed by modifying a MAXWIDE Q ME-1800 ultrasonic welding system (Minghe, Taiwan). The system includes an ultrasonic generator controller, power supply, pneumatic controller, operation time controller, stroke controller, and an automatic overload protection device.

Based on the thermal properties of XLPE (Fig.2), the conventional forming process for factory joint insulation recovery involves extrusion at 110~120∘C, followed by vulcanization molding at 160~170∘C. In the developed ultrasonic vibration-enhanced process, ultrasound is applied directly to the interface before the vulcanization step. The ultrasonic parameters were set to a frequency of 20 kHz with 50% amplitude. The vibration cycle was programmed as 4 seconds of ultrasonics followed by a 2 -second pause interval. During operation, the sonotrode was automatically traversed along the interface, as illustrated in Fig.3.

[picture]

Fig. 2. Non-isothermal DSC curve of the XLPE material used for the factory joint.

Fig. 2. Non-isothermal DSC curve of the XLPE material used for the factory joint.

[picture]

Fig. 3. Schematic diagram of the ultrasonic vibration-enhanced process for interfacial bonding.

Fig. 3. Schematic diagram of the ultrasonic vibration-enhanced process for interfacial bonding.

Real-time monitoring of acoustic pressure and temperature was performed at the interface, as shown in Fig.4. This revealed that ultrasonic vibration induces a rapid local temperature increase, reaching 130∘C within 4 seconds and stabilizing at approximately 140∘C. Furthermore, an instantaneous high-frequency acoustic pressure of about 120 kPa can be generated within milliseconds.

[picture]

Fig. 4. (a) High-frequency acoustic pressure and (b) temperature distribution curves at the interface during ultrasonic treatment.

Fig. 4. (a) High-frequency acoustic pressure and (b) temperature distribution curves at the interface during ultrasonic treatment.

For mechanical characterization, tensile tests were performed on a Sans UTM6104 universal testing machine at a crosshead speed of 200 mm/min, in accordance with the GB/T 1040-2006 standard. Thermal elongation properties of samples were evaluated following IEC 60811-202:2023 guidelines. To assess electrical performance, DC breakdown strength tests were conducted using a ZJC-100kV DC breakdown strength tester (Beijing Zhonghang Times Instrument Co., Ltd). The tests employed a pillar-to-pillar configuration with a diameter of 25 mm . A voltage ramp rate of 2kV/s was applied until breakdown occurred. Spatial charge distribution within the factory joint specimens was characterized using the pulsed electro-acoustic (PEA) method. Prior to measurement, samples were subjected to a DC electric field of 20kV/mm at 30∘C for 30 minutes (polarization phase), followed by a depolarization period of 10 minutes.

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

- Mechanical Properties Enhancement

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

The mechanical behavior of the body and recovery insulations was illustrated in Fig.5. Both maintained a tensile strength of approximately 19 MPa and an elongation at break of about 600%. These results indicated no significant difference between them. For the conventional processed interface samples, the tensile strength and elongation were measured at 10.7 MPa and 336%, respectively. With the application of ultrasonic treatment, the interfacial bonding quality was significantly enhanced, as evidenced by an increase in tensile strength to 12.7 MPa , representing an improvement of approximately 20% compared to the conventional processed samples. Additionally, the elongation at break reached 458%, corresponding to an increase of about 40%.

[picture]

Fig. 5. Tensile properties of factory joint XLPE insulation.

Fig. 5. Tensile properties of factory joint XLPE insulation.

- Thermal stability improvement

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

As illustrated in Fig.6, the thermal elongation rates of body and recovery insulation were similar, at 60% and 64%, respectively, with both exhibiting zero permanent elongation. Compared to the conventional group, ultrasonic treatment reduced thermal elongation from 127% to 110% and permanent elongation from 9% to 3%. These findings indicated that ultrasonic vibration effectively enhances the dimensional stability of the interface under thermal loading.

[picture]

Fig. 6. Thermal Elongation Properties of factory joint XLPE insulation (a) Load elongation (b) Permanent elongation.

Fig. 6. Thermal Elongation Properties of factory joint XLPE insulation (a) Load elongation (b) Permanent elongation.

- DC breakdown behavior

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

The DC breakdown strength of the insulation was characterized based on Weibull statistics analysis, as illustrated in Fig.7. The body insulation and the recovery insulation exhibited comparable breakdown levels, with characteristic breakdown strengths of approximately 395kV/mm and 360 kV/mm, respectively. In contrast, the conventional processed interface samples demonstrated a reduction in breakdown strength to about 297kV/mm. Following ultrasonic treatment, the breakdown strength increased notably to around 334kV/mm, representing an improvement of about 13% relative to the conventional processed interface samples. These findings indicated that the ultrasonic application effectively enhanced the both interfacial breakdown strength and stability of the insulation.

[picture]

Fig. 7. DC Breakdown Strength of factory joint XLPE insulation.

Fig. 7. DC Breakdown Strength of factory joint XLPE insulation.

- Space charge suppression

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

The space charge results (Fig.8) clearly demonstrated the variations in carrier injection, migration, and trapping across distinct interface conditions. During the initial phase of DC electric field application, charge injection primarily initiates near the electrode, with electrons being preferentially injected from the cathode and migrating toward the anode under the electric field influence. The charge cloud progressively extended along the sample thickness. Trap distributions and non-uniform conductivity hindered deep penetration of space charges. Consequently, charges accumulate locally.

In the conventional processed samples, interfacial negative charge accumulation was enhanced. The interface tended to retain charges, causing distortion of the local electric field and heightened energy injection, potentially triggering electrical treeing, local heat accumulation, or micro-defect discharge. Consequently, the interface became a vulnerable site for breakdown initiation, leading to reduced and more scattered breakdown strength. In contrast, ultrasonic treatment effectively suppressed the accumulation of negative charges at the interface. The high-frequency sound pressure and local temperature rise promoted secondary reinforcement of melt bonding, thereby mitigating the effects of trapping and charge retention. This resulted in a more uniform local electric field, thus enhancing breakdown strength and stability.

[picture]

Fig. 8. Space charge profiles at various time under 20kV/mm at 30∘C : (a) Charge density of conventional processed interface sample, (b) Charge density of ultrasonic processed samples, (c) Electric field of conventional processed interface sample and (d)Electric field of ultrasonic processed samples.

Fig. 8. Space charge profiles at various time under 20 k V / m m at 30 ∘ C : (a) Charge density of conventional processed interface sample, (b) Charge density of ultrasonic processed samples, (c) Electric field of conventional processed interface sample and (d)Electric field of ultrasonic processed samples.

Summary

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

This study implements ultrasonic vibration in the recovery insulation forming process in HVDC cable factory joints, achieving significant improvements in both mechanical and electrical properties at the body/recovery insulation interfaces.
(1) Through ultrasonic treatment, the interfacial tensile strength increased by approximately 20%, and the elongation at break increased by about 40%. Both thermal elongation rate and permanent elongation rate were significantly reduced, signifying significant enhancement in interfacial deformation stability and bonding quality.
(2) Compared with conventional processing, ultrasonic processing increased the DC breakdown strength of the insulation interface by approximately 13%. Furthermore, space charge accumulation was effectively suppressed, and the interfacial electric field distribution became more uniform, thereby enhancing voltage withstand capability and long-term reliability.
(3) Ultrasonic high-frequency acoustic pressure and local temperature rise effectively promoted molecular chain melting, rearrangement, and interface densification, consequently eliminating microscopic defects and charge traps.

Acknowledgment

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

This work is supported by the Science and Technology Project of State Grid Corporation of China (5500-202499389A-3-3-ZH).

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.1.pdf

G. Mazzanti, Issues and Challenges for HVDC Extruded Cable Systems, Energies 14(15) (2021) 4504. https://doi.org/10.3390/en14154504.

F.-B. Meng, X. Chen, Z. Hong, Y. Shi, H. Zhu, A. Muhammad, A. Paramane, R. Huang, Z. Han, Insulation Properties and Interfacial Quantum Chemical Analysis of Cross-Linked Polyethylene Under Different Degassing Time for HVDC Cable Factory Joint Applications, IEEE Trans. Dielectr. Electr. Insul. 30(1) (2023) 271-278. .

Y. Qin, Y. Ge, J. Zhao, Y. He, Z. Lv, K. Wu, Y. Zhao, Y. Zhou, Characteristics of Charge Accumulation at the Interface of Double-Layer XLPE Under DC Voltage, IEEE Trans. Dielectr. Electr. Insul. 32(2) (2025) 707-716. .

M. Hu, H. Zhang, X. Ge, X. Hu, Z. Zhang, X. Cao, Z. Li, W. Huo, Effect of Crosslinking Temperature on the Insulation Performance of XLPE Secondary Crosslinking Insulation Interface Layer, Polymers 17(7) (2025) 936. .

Z. Zheng, Z. Li, H. Wang, Y. Wu, B. Du, Electric Field Distribution and Evolution in Submarine Cable Factory Joint under Steady and Transient Voltages, 2024 IEEE 5th Int. Conf. on Dielectrics (ICD), 2024, pp. 1-4. .

F.-B. Meng, X. Chen, C. Dai, M. Zhang, Y. Shi, A. Paramane, A. Muhammad, Interfacial microstructure and insulation properties of 500 kV EHVDC XLPE cable factory joint under different roughness and degassing durations, Polym. Degrad. Stab. 192 (2021) 109688. .

C. Avila-Orta, C. Espinoza-González, G. Martinez, D. Bueno-Baqués, A. Maffezzoli, F. Lionetto, An Overview of Progress and Current Challenges in Ultrasonic Treatment of Polymer Melts, Adv. Polym. Technol. 32(S1) (2013) E582-E602. . 21303.

Z.A.-O. Ma, C.A.-O. Bourquard, Q.A.-O. Gao, S.A.-O. Jiang, T.A.-O. De Iure-Grimmel, R.A.O.X. Huo, X.A.-O. Li, Z.A.-O. He, Z.A.-O. Yang, G.A.-O. Yang, Y.A.-O. Wang, E.A.-O. Lam, Z.A.-O. Gao, O.A.-O. Supponen, J.A.-O. Li, Controlled tough bioadhesion mediated by ultrasound, Science 377(6607) (2022) 751-755. .

A.R.S. Santha Kumar, A. Padmakumar, U. Kalita, S. Samanta, A. Baral, N.K. Singha, M. Ashokkumar, G.G. Qiao, Ultrasonics in polymer science: applications and challenges, Prog. Mater. Sci. 136 (2023) 101113. .

V. Siciliani, R. Pelaccia, M. Alfano, D. Castagnetti, L. Raimondi, A. Zanchi, L. Donati, L. Orazi, Effect of surface preparation on adhesive bonding of CFRP compression molding laminates, Material Forming: ESAFORM 2025, 2025, pp. 2193-2202. .

Z. Zhou, C. Chen, Research advances in the ultrasonic-assisting adhesive bonding, J. Manuf. Process. 110 (2024) 134-160. .

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The original version of this eBook is available on https://www.scientific.net/book/innovative-joining-by-forming/978-3-0364-1990-9

Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 9-17

doi: 10.4028/p-5r6ON6

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2025-12-21

Revised:
2026-02-25

Accepted:
2026-03-03

Online:
2026-04-15

Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

Unni Daniel Yeniss 1,a* and Strano Matteo 1, b

¹Politecnico di Milano, Mechanical Engineering Dept., Via La Masa, 1, Milan, 20156 Italy

a* danielyeniss.unni@polimi.it, b  matteo.strano@polimi.it

*corresponding author

Keywords: joining by forming, rivet design, sheet metal joining, cold joining, low cost tooling, end of life (EoL) vehicles, sustainable manufacturing

Abstract

The direct reuse of End-of-Life (EoL) automotive sheet metal offers significant environmental benefits, particularly when panels are flattened rather than remelted. However, the geometric irregularity and variable formability of reclaimed sheets require joining solutions that operate with minimal tooling and are compatible with flat-die pressing. This work introduces a novel rivet-based joining-by-forming process designed to create a mechanical interlock using only flat tools, enabling integration into the same hydraulic press used for EoL panel flattening. Finite element simulations were performed to optimize rivet geometry and study flange formation, stress distribution, and failure mechanisms. The optimized rivet design achieves stable outward flaring under axial compression, producing a functional joint without shaped dies. Experimental tests on reclaimed automotive sheets validated the joining concept and confirmed the deformation behaviour predicted numerically. The proposed method provides a low-cost, low-energy joining strategy suitable for reclaimed steels of uncertain formability and supports the development of circular manufacturing routes based on the direct reuse of automotive sheet metal.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

The transition toward low-carbon industrial systems requires the adoption of new sustainability strategies in sectors characterized by high material and energy intensity. Transportation alone accounts for approximately 25% of global energy consumption [1] and 21.9% of world steel demand [2], with the automotive industry responsible for 19.8% of the total steel used annually [2]. In a typical 1481 kg passenger vehicle, steel constitutes nearly half of the total mass, approximately 47%, highlighting the critical role of metal use in this sector [3]. In this context, improving the resource efficiency of automotive metals represents a high-impact opportunity for advancing Circular Economy objectives beyond conventional recycling.
Farioli et al. [4] recently introduced an alternative recovery strategy for End-of-Life (EoL) automotive panels that avoids energy-intensive remelting. Their approach reshapes curved automotive body components into flat sheet metal through a controlled flattening operation using a hydraulic press. The overall recovery workflow consists of four primary stages (Fig. 1):
(1) selection of EoL vehicle body panels;
(2) material extraction through cutting and disassembly;
(3) flattening of the curved panels into planar sheets; and
(4) paint removal using mechanical, chemical, or laser-assisted methods.

[picture]

Fig. 1. Scheme of the EoL sheet metal recovery process.

Fig. 1. Scheme of the EoL sheet metal recovery process.

This process offers substantial environmental benefits compared to recycling. For each kilogram of steel recovered through flattening, approximately 2.2 kg of CO2 emissions and 24 MJ of energy can be saved [4]. Additional evidence from Ali et al. [5] shows that direct reuse of scrap sheet metal within manufacturing can reduce costs by 40% and energy consumption by 67%. They demonstrated that irregularly shaped automotive scrap sheets can be directly reused by employing a Voronoi-based geometric algorithm to generate façade panels from heterogeneous offcuts. Since the scrap pieces differ significantly in shape and size, the Voronoi tessellation allows the target surface to be subdivided into polygons that can be efficiently matched to the available sheets, minimizing trimming waste and avoiding remelting.
This example highlights a crucial constraint of upcycling workflows: reclaimed sheets must often be combined to obtain larger functional components, making an efficient, low-tooling joining method essential.
The geometrical and material variability of End-of-Life (EoL) automotive panels introduces significant technical challenges for their direct reuse. Recovered sheets typically exhibit heterogeneous dimensions, variable thickness, pre-existing damage, surface degradation, and reduced or uncertain formability due to ageing and prior deformation history. Consequently, the reuse of flattened EoL sheet metal requires a rethinking of downstream manufacturing operations. In particular, the intrinsically limited size of individual body panels makes joining operations essential to produce larger sheets or structural subassemblies suitable for secondary applications.
Available joining-by-forming technologies such as clinching, self-piercing riveting (SPR), and double-sided SPR are widely recognized for providing reliable mechanical joints with comparatively low environmental impact relative to fusion welding [6, 7]. However, these processes rely on shaped dies, controlled material ductility, and dedicated tooling. Such requirements are difficult to satisfy when working with reclaimed automotive sheet metal characterized by thickness variability, unknown forming history, and degraded or uncertain ductility. Blind riveting appears to be a suitable solution, but it suffers from intrinsic drawbacks, including low process speed and the formation of protrusions incompatible with functional or aesthetic surface requirements. More generally, existing rivet-based solutions mostly require pre-formed features, contoured dies, or proprietary equipment, which limits their integration into flexible, low-cost reuse workflows.

Novelty of the present Work.

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

The contribution of this study is a rivet-based joining-by-forming concept explicitly engineered for flattened End-of-Life (EoL) automotive sheets, where thickness, surface condition, and ductility are often uncertain and where the available equipment is typically a flat-tool hydraulic press used for panel flattening. Unlike clinching and self-piercing riveting, which rely on shaped die cavities and controlled sheet flow, the proposed joint is generated with two flat dies: a purposely profiled rivet stem is axially compressed to trigger a repeatable outward buckling mode, forming a circumferential flange that mechanically locks the sheet stack. The rivet geometry (chamfer, stem length, and headshank radius) is designed to (i) stabilize the buckling path, (ii) promote outward flaring rather than inward collapse, and (iii) tolerate high pressing forces without dedicated tooling. This enables the joining operation to be performed in the same press and with the same flat tooling already used for EoL panel flattening, reducing capital cost and simplifying integration into reuse workflows. To the authors' knowledge, this is the first rivet-based joining-by-forming approach demonstrated to create a functional mechanical interlock under flat-die conditions and validated numerically and experimentally on reclaimed automotive sheet metal.
In Fig. 2, two examples are shown of upcycled components produced from reclaimed automotive panels, demonstrating how geometrically limited and irregular sheets can nonetheless be transformed into high-value products, underscoring the need for efficient joining solutions to enable such applications.

[picture]

Fig. 2. Examples of products obtained from upcycled automotive sheet metal. From left to right, Tables obtained from scrapped automotive sheets, made by Weld House (Photo by Dornob). A table obtained from an Opel Admiral car, made by Unibro Design.

Fig. 2. Examples of products obtained from upcycled automotive sheet metal. From left to right, Tables obtained from scrapped automotive sheets, made by Weld House (Photo by Dornob). A table obtained from an Opel Admiral car, made by Unibro Design.

Designed Rivet Working Principle

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

The proposed joining concept is designed to operate using only a flat punch and flat die, identical to the tools already employed in the EoL sheet-flattening process. This allows the joining operation to be seamlessly integrated into the same hydraulic pressing workflow without additional tooling. The joint is formed through a sequence of deformation stages in which a specially designed rivet undergoes controlled buckling and flange formation to create a mechanical interlock between the sheets.
The process begins with the insertion of the rivet into two pre-drilled holes in the sheet stack, with a clearance of 0.1 mm between the rivet shank and the hole diameter (Fig. 3a). The rivet head is positioned on the die side, while the stem faces the punch. As the punch advances, the rivet stem is compressed axially and undergoes progressive buckling, causing the lower portion of the stem to flare outward and form an expanding flange (Fig. 3b). Concurrently, the sheets are forced into contact with the underside of the rivet head, and slight upward bending of the upper sheet may occur as the stack accommodates the developing flange (Fig. 3c).
Continued punch displacement forces the flanged region into a fully flattened configuration, securing the sheets between the rivet head and the newly formed flange (Fig. 3d). The resulting interlock is generated entirely through plastic deformation of the rivet stem and the localized bending of the sheets, without relying on a shaped die cavity. This characteristic differentiates the proposed solution from conventional riveting, clinching, and self-piercing riveting processes, and enables a low-cost joining operation compatible with reclaimed sheet metal of variable formability.

[picture]

Fig. 3. Representation of the proposed rivet joining method using flat tools.

Fig. 3. Representation of the proposed rivet joining method using flat tools.

Numerical and experimental investigations (described in the following sections) demonstrate that rivet geometry, particularly the chamfer length, head-shank radius, and stem length plays a decisive role in controlling the buckling mode, flange extent, and final interlock quality. The deformation path must be carefully tuned to ensure outward flaring rather than uncontrolled buckling, while minimizing excessive sheet displacement or surface irregularities.

Both numerical simulations, using Transvalor Forge, and experimental tests, with a Pezzato CB AUTO hydraulic press machine, have been performed on the novel rivet design. Experimental tensile tests have been performed with an MTS Alliance RT/100 machine. Both numerical simulations and experimental tests were conducted using DC04 low carbon steel sheets (Table 1). Stainless Steel 303, sourced from the Xometry materials catalog, was selected as the rivet material. For the experimental phase, rivet prototypes were CNC machined from this material.

Table 1. Materials used in the tests. DC04 steel characteristics are based on Trzepieciński and Najm work [8].

Grade | Yield Stress [Mpa] | Tensile Strength [Mpa] | Elongation [%]
Stainless Steel 303 | 351 | 398 | 31
DC04 Steel | 210 | 270 - 350 | 38

Numerical Simulations

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

Finite element simulations were conducted using Transvalor Forge to evaluate the deformation behavior of the proposed rivet geometry and to investigate flange formation, sheet displacement, and failure tendencies under flat-die compression. All simulations employed 1 mm thick DC04 steel sheets and were performed with a pressing strength of 25 tons. The rivet material was modeled using stainless steel 303 properties, as reported in Table 2.
Due to the axisymmetric nature of the initial rivet deformation process, most simulations were performed in a 2D axisymmetric configuration, enabling reduced computational cost while accurately capturing the buckling and flanging behavior of the rivet stem. The punch and die were modeled as rigid bodies, whereas the rivet and sheets were modeled as elasto-plastic deformable solids.
Material behavior was described using an isotropic hardening law, calibrated from tensile tests for the reclaimed DC04 sheets and literature data for the AISI 303. A Coulomb-limited Tresca friction model was applied at all contact interfaces, using a Coulomb friction coefficient of 0.05 and a Tresca shear factor of 0.1 , to represent a low-friction metal-to-metal interaction. The Coulomb value ( μ= 0.05 ) is commonly adopted in sheet-metal forming to model low friction, while the Tresca parameter ( m=0.1 ) was set in accordance with the software guidelines, which recommend using a Tresca value approximately twice the Coulomb coefficient. Although the selected coefficients provide a reasonable and stable baseline for the present study, a dedicated sensitivity analysis is currently underway to select the optimal friction coefficients. The adopted values should be regarded as starting minimal friction coefficients. Mesh refinement was applied in regions of high expected strain gradients, with a minimum 2D element size of approximately 0.1 mm , ensuring accurate capture of the buckling mode and localized thinning. Typical computation times ranged from approximately 30 to 50 minutes for 2D simulations to more than 5 hours for full 3D models. Numerical simulations were validated by visually comparing the simulations results and the pressed rivet in the experimental phase, which will be discussed more in detail later in this work.
The rivet geometry (Fig. 4) is inspired by commercially available self-piercing rivets and incorporates the following distinctive features:

A 4.0 mm stem length and 3.3 mm diameter, ensuring sufficient material flow and outward flaring under compression to create a stable mechanical interlock.

An extended, sharpened chamfer designed to minimize the contact area between the flat surface at the end of the chamfer and the hydraulic press punch. This geometry enables earlier
contact between the punch and the chamfer, thereby reducing the risk of inward buckling, which would lead to joint failure, and promoting controlled outward flaring necessary for joint formation.

A reduced fillet radius between the underside of the rivet head and the base of the stem, maximizing the contact area between the rivet and the sheet metal during pressing and enhancing load transfer within the joint.
The rivet when pressed, correctly flanged outwards leading to the creation of a constraint in the punch side and on the die side thanks to the sheets being pushed against the head, granting further resistance to the joint (Fig. 5).

[picture]

Fig. 4. Section view of the new version of the rivet.

Fig. 4. Section view of the new version of the rivet.

[picture]

Fig. 5. Result of the numerical simulation on the third tested rivet. A working connection is formed.

Fig. 5. Result of the numerical simulation on the third tested rivet. A working connection is formed.

The numerical simulations allowed to see the Latham Cockroft normalized criterion on the pressed rivet and assess its damage mode when the joint is being created. Both in the initial setup and reversed orientation, the rivet showed high stresses and fractures around the border of the flared stem (Fig. 6), leading to cracks that are radially distributed similar to an asterisk. Numerical simulations showed that in the rivet with the head lying on the die side, like in fig. 3, presented deeper but more localized fractures (Fig. 6a); whereas the rivet in reversed orientation, with the head pointing towards the punch side, showed less deep but more numerous fractures leading to a less defined and more wavy border of the pressed stem, compared to the rivet pressed in with the head pointing towards the die side (Fig.

6b). A reason of this difference in behaviour can be found in the fact that the rivet when pressed in its reversed orientation, has its stem inserted through the pre-drilled holes and the material flow when pressed is constrained by the material of the parts to join, whereas the rivet when pressed with the head pointing towards the die side has the stem bulging out of the pre-drilled hole and finds less constraints in its flow when pressed.

[picture]

Fig. 6. Latham Cockroft normalized Criterion of the pressed rivet (a) and of the pressed rivet with reversed orientation (b). The simulation on the reversed rivet (b) is symmetrical and was performed only on one half of the model.

Fig. 6. Latham Cockroft normalized Criterion of the pressed rivet (a) and of the pressed rivet with reversed orientation (b). The simulation on the reversed rivet (b) is symmetrical and was performed only on one half of the model.

To estimate joint strength, push-out simulations were performed in which a punch applied force through the hollow rivet core while the sheets were clamped between blankholders (Fig. 7). Pushout simulations showed an estimated maximum force of the joint of 2.503 kN .

[picture]

Fig. 7. Setup of the pushout simulations. Red and yellow are the blankholders, the punch is the brown element. The dark green lower die has a hole to allow the falling of the pushed out rivet.

Fig. 7. Setup of the pushout simulations. Red and yellow are the blankholders, the punch is the brown element. The dark green lower die has a hole to allow the falling of the pushed out rivet.

[picture]

Fig. 8. Plots of the required force vs. pressing time in the pushout numerical simulations.

Fig. 8. Plots of the required force vs. pressing time in the pushout numerical simulations.

Experimental Riveting Tests

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Experimental tests were conducted to validate the feasibility of the rivet-based joining process and to compare the deformation behavior observed in simulations with real pressed samples. Prototypes of the optimized rivet were CNC-machined from AISI 303 stainless steel, matching the geometry and material properties used in the numerical models. All experiments were performed using a Pezzato CB AUTO hydraulic press, equipped with flat punch and die surfaces identical to those used for the EoL sheet-flattening process and set on 25 tons pressing force.
Automotive sheet-metal offcuts recovered from hoods and roofs were selected as representative EoL materials, and a sample of 0.99 mm thickness was selected, comparable to the 1.0 mm used in the FEM simulations, enabling direct comparison between numerical and experimental results. Each dogbone specimen was cut at mid-length, and a single 3.5 mm diameter hole was drilled 6.5 mm from the cut edge. The rivet was inserted into this hole before pressing.
During pressing, the rivet stem underwent significant plastic deformation, forming the outward flange responsible for the mechanical interlock. In both orientations, fractures developed along the deformed stem, consistent with high strain localization predicted by simulations.
Fig. 9 illustrates representative pressed rivets used to join samples obtained from EoL automotive sheet metal. The rivet flared outward to form a retaining flange. The upper sheet exhibited slight upward bending, in agreement with simulation predictions. No fractures were observed in the sheet metal itself. The pressed rivet on the left hand side of the picture was pressed with the rivet head lying on the die side and the stem pointing towards the punch, whereas the right hand side joint presents the rivet reversed with the head pointing towards the punch side. In both cases the rivets formed reliable interlocks and confirm that the geometry is robust and that flange formation is reliably achieved using flat dies, without requiring any contoured or specialized tooling. The fractures starting from the borders of the rivet and disposed circularly were expected from the numerical simulations results and are caused by the high pressing force. The fractures around the rivets are caused by the pressing force and affect only the paint layer over the EoL sheet metal from the samples, the sheets material was not damaged.

[picture]

Fig. 9. On the left hand side the pressed rivet with the initial setup orientation, on the right hand side the rivet pressed in the reversed direction with the head on the punch side.

Fig. 9. On the left hand side the pressed rivet with the initial setup orientation, on the right hand side the rivet pressed in the reversed direction with the head on the punch side.

16 tensile tests were performed on the pressed samples using a MTS Alliance RT/ 100 universal testing machine. Both single rivets and double rivets setups were tested in both orientations, as shown in Fig. 3 and with the setup reversed, where rivet head points towards the punch side. The reversed rivets showed a slightly lower load-bearing capacity.

Table 2. Maximum loads of the tensile tests of single rivet setups. (r) indicates a rivet pressed with reversed orientation.

Single Rivet Setup
Thickness (mm) | Test 1 mathematical formula | Test 2 mathematical formula | Test 3 mathematical formula | Average mathematical formula
0.75 | 1.17 | 1.499 | 1.174 | 1.281
0.83 | 1.297 | 1.449 | superscript number | 1.33
0.99 | 1.785 | 1.624 | superscript number | 1.71

Table 3. Maximum loads of the tensile tests of double rivet setups. (r) indicates a rivet pressed with reversed orientation.

Double Rivet Setup
Thickness (mm) | Test 1 (kN) | Test 2 (kN) | Test 3 (kN) | Average (kN)
0.75 | 2.306 | 2.12 | - | 2.213
0.83 | 2.704 | 2.646 | - | 2.675
0.99 | 2.464 | 2.552 | superscript number | 2.524

Validation of Numerical Simulations.

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The numerical model was validated by comparison with experimental observations in terms of deformation mode and damage localization. The simulations predict a progressive axial compression of the rivet stem followed by a controlled outward buckling and flaring, producing a mechanical interlock under flat-die conditions. The experimental tests, performed by applying a 25 ton pressing force to the joint, exhibited the same deformation sequence, confirming that the model captures the governing mechanisms and reproduces the distinct response observed for the standard and reversed orientations. In addition, simulations evaluated with the normalized Latham-Cockcroft criterion indicate damage localization at the perimeter of the flared rivet stem. Consistently, experiments revealed radial cracks initiating at the same location, while no fractures were observed in the sheet material, as predicted. The only additional cracking observed in the specimens occurred in the paint coating of the painted EoL samples, which showed superficial coating cracks without substrate failure.

Conclusion

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This study demonstrated the feasibility of a rivet-based joining-by-forming process specifically designed for flat-die operation, enabling direct integration into EoL sheet-metal flattening workflows. Numerical simulations highlighted the critical influence of rivet geometry on buckling and flange formation, leading to an optimized design capable of forming a stable mechanical interlock without dedicated dies. The proposed approach addresses the geometric variability and uncertain formability of reclaimed automotive sheets by eliminating the need for shaped dies and dedicated joining equipment.
Finite element simulations were used to analyse rivet deformation, flange formation, and damage evolution under axial compression. The results demonstrated that rivet geometry particularly stem length and chamfer configuration governs the buckling mode and the stability of outward flaring. The proposed rivet promoted controlled radial outward expansion of the stem, resulting in a stable mechanical interlock while limiting sheet deformation. Damage predictions based on the normalized Latham-Cockcroft Normalized Criterion indicated localized cracking at the flared rim of the rivet, without critical failure of the sheets or of the joint. Numerical push-out simulations on a joint composed of a rivet estimated a maximum load of approximately 2.5 kN .
Experimental tests confirmed the numerical simulations with the creation of a stable interlock in both orientations, with no structural damage on the sheets found. Fractures were confined to the pressed rivet stem border and the paint coating layer on samples that were painted. Tensile tests yielded average maximum load of 1.71 kN for single rivet configurations joining 0.99 mm thick samples and 2.675 kN for double rivet configurations joining 0.83 mm thick samples.

Overall, the results demonstrate that mechanically reliable joints can be achieved using exclusively flat dies, eliminating the need for shaped cavities or dedicated joining systems. The proposed process provides a low-tooling, low-energy solution suitable for circular manufacturing routes based on direct reuse of automotive sheet metal.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.9.pdf

«International Energy Outlook 2016 With Projections to 2040», U.S. Energy Information Administration, mag. 2016. doi: https://doi.org/10.2172/1296780.

«European Steel in Figures - Covering 2023», Eurofer, 2024. .

F. E. K. Sato e T. Nakata, «Energy consumption analysis for vehicle production through a material flow approach», Energies, vol. 13, fasc. 9, mag. 2020, doi: 10.3390/en13092396.

D. Farioli, M. Fabrizio, E. Kaya, V. Mussi, e M. Strano, «Energy measurements and LCA of remanufactured automotive steel sheets», in Materials Research Proceedings, Association of American Publishers, 2023, pp. 1947-1956. doi: 10.21741/9781644902479-210.

A. K. Ali, Y. Wang, e J. L. Alvarado, «Facilitating industrial symbiosis to achieve circular economy using value-added by design: A case study in transforming the automobile industry sheet metal waste-flow into Voronoi facade systems», Journal of Cleaner Production, vol. 234, pp. 1033-1044, ott. 2019, doi: 10.1016/j.jclepro.2019.06.202.

G. Meschut et al., «Review on mechanical joining by plastic deformation», Journal of Advanced Joining Processes, vol. 5. Elsevier B.V., june 2022. doi: 10.1016/j.jajp.2022.100113.

J. Wu, C. Chen, Y. Ouyang, D. Qin, e H. Li, «Recent development of the novel riveting processes», doi: 10.1007/s00170-021-07689-w/Published.

T. Trzepieciński e S. M. Najm, «Current Trends in Metallic Materials for Body Panels and Structural Members Used in the Automotive Industry», Materials, vol. 17, no. 3, p. 590, jan. 2024, doi: 10.3390 /ma17030590.

This ePub file was created using machine-generated methods and automated software to convert the source text into an ePub format and may contain formatting inconsistencies from the original version. Formatting may vary depending on the device used.

The original version of this eBook is available on https://www.scientific.net/book/innovative-joining-by-forming/978-3-0364-1990-9

Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 19-28

doi: 10.4028/p-aok1H1

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2025-12-21

Revised:
2026-02-13

Accepted:
2026-02-23

Online:
2026-04-15

Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

The original version of this paper is available on https://www.scientific.net/MSF.1185.19.pdf

Jessica Sarris 1,a*, Michael Lechner 1, b

¹Institute of Manufacturing Technology, Friedrich-Alexander-Universität Erlangen-Nürnberg, Egerlandstr. 13, 91058 Erlangen, Germany

ajessica.sarris@fau.de, b  michael.lechner@fau.de

Keywords: Self-Piercing Riveting (SPR), Versatile Joining, Asymmetric Tumbling, Tumbling SelfPiercing Riveting (T-SPR)

Abstract

Conventional self-piercing riveting (SPR) produces rotationally symmetric joints with largely uniform mechanical behavior. While this provides robust performance in many applications, increasing demands for material-efficient lightweight design, complex load paths and hybrid material systems require more versatile joining strategies. Recent experiments have demonstrated that an adapted tumbling SPR (T-SPR) process can intentionally induce non-rotationally symmetric joint geometries and thereby extend process limits beyond those of conventional SPR. Such asymmetric joints offer the potential to tailor load-bearing capacity and energy absorption to specific load directions, which could be particularly advantageous in crash-relevant or multi-material applications. Building on these findings, the present study shifts the focus from geometry control to the systematic evaluation of the mechanical performance of asymmetric T-SPR joints. Specimens were produced using T-SPR with tailored combinations of tumbling angle and velocity. The joints are manufactured with a versatile tumbling self-piercing riveting tool. To assess the resulting mechanical properties, cross tensile and tensile shear tests are conducted. From the resulting force-displacement curves, typical mechanical properties such as ultimate load, load-bearing capacity, displacement at failure and absorbed energy are derived. The mechanical performance of asymmetric joints is evaluated in comparison with symmetric reference joints produced by tumbling self-piercing riveting. This enables both the demonstration of direction-dependent mechanical behavior of asymmetric joints compared to symmetric references and a systematic evaluation of how geometric anisotropy affects load-bearing capacity, absorbed energy and failure characteristics.

Introduction

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Lightweight design is one of the key elements for achieving the targets of reducing CO2 emissions over the entire vehicle life cycle [1]. In this context, application-specific material design is of central importance [2]. The industrial implementation of such lightweight structures requires joining technologies that are capable of connecting materials with high process reliability and without thermal influence on the microstructure [3]. Mechanical joining, in particular SPR, has become established as the dominant technology in the automotive industry [4]. In addition to the high process forces required and the need for a rigid machine setup [5], the rigid and purely axial punch kinematics produce rotationally symmetric joints. SPR therefore generates direction-independent joint properties, although body-in-white structures are often subjected to complex, anisotropic load paths, for which a directionally tailored adaptation of joint properties would be advantageous. This is why adaptable joining processes are required. One approach extending the process limits is tumbling self-piercing riveting (see Fig. 1). T-SPR is an incremental further development of the conventional SPR, in which the completely translational axial motion of the punch is superimposed by a tumbling kinematics [6]. In contrast to conventional self-piercing riveting, the punch in T-SPR has a conical shape and therefore does not act on the rivet over its entire contact area [7].

[picture]

Fig. 1. Tumbling self-piercing riveting (T-SPR) process and tumbling kinematics

Fig. 1. Tumbling self-piercing riveting (T-SPR) process and tumbling kinematics

The conical tumbling punch performs a rolling motion on the rivet head, and the resulting reduced contact area leads to lower maximum process forces [8]. In contrast to SPR, the forming force is not introduced globally but is concentrated locally [9]. The individual process steps of T-SPR, namely clamping, piercing, spreading and setting, are identical to those of conventional self-piercing riveting [10], with the exception of the tumbling punch. In this study, the load-bearing capacities and fracture behavior of asymmetrically tumbled self-piercing riveted joints are investigated in comparison with symmetrically tumbled reference joints with two identical joining partners.

Experimental Test Setup

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The material used in this study is a precipitation-hardenable aluminum alloy of the 6xxx series on an Al-Mg-Si basis, EN AW-6014 in the T4 temper, with an initial sheet thickness of to =2.0 mm on the die side and a sheet thickness of t0=1.0 mm on the punch side. The alloy has been developed for automotive outer panel applications and is characterized by a high forming capability [11].

Joining by tumbling self-piercing riveting (T-SPR). The joining experiments have been realized on a universal testing machine (type Walter + Bai FS-300) equipped with an integrated versatile tumbling tool (see Fig. 2) [6], which simultaneously records the force-displacement curves. By integrating the versatile tool into the universal testing machine, the punch stroke is directly applied via the machine's crosshead movement in the z -direction. The tumbling motion is generated by a twoaxis drive mechanism consisting of a linear and a rotary actuator. The kinematic coupling of the two axes enables controlled positioning at any point on the circular path and thus freely programmable tumbling strategies.

[picture]

Fig. 2. Tumbling self-piercing riveting tool and tool specifications

Fig. 2. Tumbling self-piercing riveting tool and tool specifications

Within the scope of this study, both symmetric and asymmetric tumbling strategies are considered. For the symmetric strategies, tumbling angles of α=1∘ and α=5∘ are investigated. In the symmetric 1∘ strategy, the punch is moved around the tool axis at a constant tumbling angle of α=1∘ starting from the tumbling onset ht0. The angle remains constant over the entire tumbling phase and is reduced to zero during one final tumbling revolution. The symmetric 5∘ strategy is implemented equally. The asymmetric 1∘/5∘ strategy is defined such that, within one full revolution of the tumbling punch, 180∘ of the rivet circumference are tumbled with α=1∘ and the remaining 180∘ with α=5∘. This angle profile is repeated continuously until the end of the tumbling process. The onset of tumbling is chosen such that the rivet has fully pierced the upper sheet layer before the tumbling motion is activated, in order to prevent rivet tilting. Tumbling velocities of φ=90∘/s,180∘/s and 360∘/s are investigated. The crosshead velocity Vt is adjusted in each case to keep the stroke per revolution constant, thereby changing only the temporal sequence of the local forming steps and, consequently, the local forming speed. This makes it possible to separate the influence of tumbling velocity from geometric effects. The tool is equipped with a blank holder, which clamps the sheets against the die during joining and thus prevents bulging or lateral displacement of the sheets. The process parameters are summarized in Table 1. For each parameter combination, n=3 repetitions are carried out.

Table 1. Parameters of the investigation

Die
Rivet

Tumbling

angle mathematical formula

Tumbling onset

Tumbling

velocity mathematical formula

Tumbling

kinematic

FM 090 2016 | mathematical formula | superscript reference | 2.5 | superscript reference | Circle

Testing of Load-Bearing Capacity

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To investigate the load-bearing capacity of the asymmetrically tumbled self-piercing riveted joints, mechanical tensile shear and cross tension tests are carried out on the specimens. The tests are performed on a ZwickRoell Z10 universal testing machine with a maximum force of 10 kN under quasi-static loading in accordance with [12]. All tests are conducted at a constant crosshead velocity vt=10 mm/min.

Tensile shear tests. To investigate the load-bearing capacity of the joined connections, the tensile shear strength of specimens with dimensions of 45×105 mm2 is determined. The focus of this investigation is on the orientation of the asymmetry. Therefore, three different orientations of the maximum tilting ( 1∘ side) are considered (Fig. 3). The different orientations of the tilting are defined as follows: (a) maximum tilting oriented towards the 2.0 mm lower sheet, (b) maximum tilting oriented towards the 1.0 mm upper sheet, and (c) maximum tilting located on the left-hand side of the loading direction.

[picture]

Fig. 3. Illustration of the tensile shear test and the orientations (a), (b) and (c)

Fig. 3. Illustration of the tensile shear test and the orientations (a), (b) and (c)

The joining partners are joined with an overlap length of 16 mm . The specimen is clamped in the two grips of the universal testing machine with a constant clamping force (Fig. 3). The free clamping
length is 95 mm , according to [12]. During the tests, force and displacement are recorded. In addition, the tensile shear tests are monitored using the Aramis (GOM) digital image correlation system in order to capture displacement and strain fields on the specimen surface, to determine local strains and to analyze the failure behavior. For this purpose, the punch-side surfaces in and around the overlap region are coated with a high-contrast speckle pattern.
Cross tension tests. In addition to the tensile shear tests, cross tension tests are carried out. For this purpose, specimens according to DVS/EFB 3480-1 [12] with dimensions of 50×150 mm2 are used (Fig. 4). The riveted joint is positioned centrally within the overlap of the sheets. The joined specimens are clamped in two specimen holders that are fixed to the upper and lower fixtures of the universal testing machine. During the test, the force and displacement data are recorded by the testing machine. Two orientations of the tilting ( 1∘ side) were investigated: in the first case (a), the tilting is aligned with the longitudinal direction of the 1 mm upper sheet, corresponding to a rotation around an axis parallel to the sheet width 1w, whereas in the second case (b) the tilting is rotated by 90∘ relative to this configuration and is thus aligned with the sheet width 1w, corresponding to a rotation around an axis parallel to the sheet length 1 .

[picture]

Fig. 4. Illustration of the cross tension test and the orientations (a) and (b)

Fig. 4. Illustration of the cross tension test and the orientations (a) and (b)

Results and Discussion

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Tensile shear tests. The tensile shear tests are used to evaluate the joint strength of the specimens. To quantify the influence of the process parameters and the geometrical anisotropy, the mean maximum tensile shear forces ( F―max  ) a) are evaluated. In addition to the maximum force, the energy absorption b) of the joint is a decisive indicator of the failure behavior. As an evaluation criterion, the dissipated work up to a decrease of the tensile shear force to 30% of its maximum value ( W0,3 Fmax  ) is considered [12]. Fig. 5 shows the mechanical characteristics as a function of tumbling velocity, orientation of the maximum tilting and the symmetric and asymmetric tumbling strategies.

[picture]

Fig. 5. Mean maximum tensile shear forces and work of specimens for the different tumbling velocities and strategies (symmetric and asymmetric) as a function of orientations (a), (b) and (c)

Fig. 5. Mean maximum tensile shear forces and work of specimens for the different tumbling velocities and strategies (symmetric and asymmetric) as a function of orientations (a), (b) and (c)

The symmetric 1∘ reference specimens reach mean tensile shear forces between approximately 1.9 and 2.03 kN , whereas the symmetric 5∘ specimens reach clearly higher values of about 2.2−2.3kN, when considering the forces across all investigated tumbling velocities. Thus, a larger tumbling angle increases the tensile shear load-bearing capacity by around 18%. In comparison, the influence of the tumbling velocity φ is small. The average increase in force of about 17.5% confirms that the larger tumbling angle of α=5∘ generates a more noticeable undercut due to stronger radial material displacement, which provides a higher resistance against tensile shear loading. The 1∘ reference group exhibits the lowest energy absorption capacity with approximately 7.6−9.2 J, whereas the 5∘ specimens reach significantly higher values of about 10 J . This correlates with the more pronounced undercut of the 5∘ joints.
For the asymmetrically tumbled joints (α=1∘/5∘), tensile shear forces between 2.15 kN and 2.32 kN and dissipated energies up to 11.2 J are achieved. On average, these values lie slightly above the symmetric references ( F―max, symm =2.14kN ), representing the combined mean of the symmetric 1∘ and 5∘ reference specimens. In particular, orientation (a) (maximum tilting towards the 2 mm lower sheet) reaches or even exceeds the level of the symmetric 5∘ references, which indicates an undercut that is more effective in the shear direction. As shown in previous investigations, the undercut on the 1∘ tumbled side is more pronounced than on the 5∘ side. The larger angle in the 5∘ sector results in higher local contact pressure and a stronger geometric deformation of the rivet head by the punch [13]. Due to the larger contact area, the opposite 1∘ sector is subjected to lower local forming pressure. As a consequence, the rivet and sheet material displaced on the 5∘ side may flow into the 1∘ sector, causing the rivet foot to flare more strongly in the radial direction and thereby increasing the undercut. A dependence of the shear strengths on the loading rates for the asymmetric specimens is hardly noticeable. It can be seen that orientation (a) reaches the highest force at φ=180∘/s, whereas (b) and (c) show similar values across the three velocities. For the symmetric specimens, a stronger dependence on the velocities is observed. As φ increases, both the forces and work increase.

Failure behavior in tensile shear tests. In the following, the failure behavior of the individual orientation groups (a), (b) and (c) is analyzed. Fig. 6 shows the tensile shear forces of the different
orientation groups. All joints were produced using the asymmetric tumbling strategy with a tumbling angle of α=1∘/5∘ and a tumbling velocity of φ=180∘/s.

[picture]

Fig. 6. Force-displacement curves of the asymmetrically tumbled joints for the three orientations (a), (b) and (c) under tensile shear forces

Fig. 6. Force-displacement curves of the asymmetrically tumbled joints for the three orientations (a), (b) and (c) under tensile shear forces

At the beginning of the test, the force increases similarly for all specimens up to a displacement of approximately x=0.5 mm. Subsequently, the force-displacement curves diverge depending on the orientation of the maximum tilting. Notably is specimen a_3, which exhibits the highest tensile shear maximum and the largest displacement. For interpretation, the major and minor strains as well as the tensile strains in loading direction cy are analyzed for the three orientations at displacements of x=1.0−4.0 mm and a tumbling velocity of φ=180% s (Fig. 7, Fig. 8 and Fig. 9).

[picture]

Fig. 7. Comparison of major strains for orientations (a), (b) and (c)

Fig. 7. Comparison of major strains for orientations (a), (b) and (c)

Fig. 7 shows the evolution of the major strain fields for orientations (a), (b) and (c). For orientation (a), the major strains localize in a narrow band below the rivet along the shear path, which is consistent with the highest tensile shear force and energy absorption. In contrast, orientations (b) and (c) exhibit more distributed and partly laterally shifted major strain zones, indicating a less efficient utilization of the local deformation capacity in the main load path.

[picture]

Fig. 8.Comparison of minor strains for orientations (a), (b) and (c)

Fig. 8. Comparison of minor strains for orientations (a), (b) and (c)

Fig. 8 illustrates the corresponding minor strain fields. When looking at the strains, a continuous expansion of the blue compressive zone can be observed between 1 and 3 mm . This is attributed to the material increasingly contracting in the rivet region to compensate for the longitudinal strain. At 4 mm displacement, interruptions of the compressive zone appear for (b) and (c) in the form of compression (red). Orientation (a) exhibits a more ductile and stable response, with a smooth transverse contraction up to failure. In contrast (b) and (c) show a more unstable end-stage behavior, characterized by localized minor-strain hotspots at x=4 mm, indicating pronounced, highly concentrated constriction or shear deformation that is not observed for orientation (a).

[picture]

Fig. 9. Comparison of strains in y-direction (loading direction) for orientations (a), (b) and (c)

Fig. 9. Comparison of strains in y-direction (loading direction) for orientations (a), (b) and (c)

Fig. 9 compares the strains in y-direction (loading direction) for the three orientations. Orientation (a) is characterized by smaller regions of tensile strain (red) and more extensive compressive zones (blue) compared to orientations (b) and (c). At x=4 mm, orientation (a) shows a more diffuse and slightly wider distribution of tensile strains (red). In contrast, for orientations (b) and (c), the red regions are more pronounced and concentrated around the rivet, indicating a different deformation behavior compared to (a). One possible reason is the arrangement of the tumbling angles, which leads to a direction-dependent load path in the rib. Depending on the orientation, the tensile forces are either distributed over a wider area (orientation (a)) or introduced more locally into the material (orientations (b) and (c)). This interpretation is consistent with the tensile shear results in Fig. 5, where orientation (a) exhibits the highest tensile shear forces.

Cross tension tests. Fig. 10 shows the mean maximum cross tension forces a) and the dissipated work b) up to a decrease of the cross tension force to 30% of its maximum value (W0,3 Fmax) [12] as
a function of tumbling velocity, tumbling strategy (symmetric/asymmetric) and orientation of the maximum rivet tilting (a) and (b).

[picture]

Fig. 10. Mean maximum cross tension forces and dissipated work of the specimens for the different tumbling velocities and tumbling strategies (symmetric and asymmetric) as a function of orientations (a), (b) and (c)

Fig. 10. Mean maximum cross tension forces and dissipated work of the specimens for the different tumbling velocities and tumbling strategies (symmetric and asymmetric) as a function of orientations (a), (b) and (c)

The cross tension parameters indicate an overall robust process behavior. The mean maximum cross tension forces for all investigated parameter combinations lie within a narrow range of approximately 1.67−1.74kN. Neither the tumbling strategy (symmetric vs. asymmetric) nor the tumbling angle or tumbling velocity exhibits a pronounced influence. The work is about 10 J for all groups, with the 5∘ references reaching only slightly higher values than the 1∘ specimens. The asymmetrically tumbled joints thus achieve, in cross tension, the same level as the symmetric references and tend to exhibit slightly lower scatter.

Failure behavior in cross tension tests. The failure patterns of all specimens are very similar. In all tests, the rivet remains in the matrix-side lower sheet (t0=2.0 mm), while the punch-side upper sheet is pulled out of the rivet interlock (cross tension failure by sheet pull-out). No systematic differences in failure behavior could be identified between the symmetrically and asymmetrically tumbled joints.

Summary and Outlook

The original version of this paper is available on https://www.scientific.net/MSF.1185.19.pdf

This study investigates the load-bearing capacities of asymmetrically tumbled self-piercing riveted joints produced by a tumbling self-piercing riveting process in aluminum-aluminum (EN AW-6014 T4) sheet combinations with a die-side sheet thickness of t0=2.0 mm and a punch-side sheet thickness of t0=1.0 mm. Tensile shear tests, including digital image correlation using Aramis (GOM), as well as cross tension tests were carried out in order to evaluate the load-bearing capacities of the asymmetric riveted joints in comparison with symmetric reference joints. In addition, the failure behavior was analyzed.
The key findings can be summarized as follows:

The asymmetric tumbling strategy ( α=1∘/5∘ ) enables a targeted adjustment of anisotropic strength properties.

The asymmetric joints do not reach the mean value of the symmetric references but reach the tensile shear load-bearing level of the 5∘ reference and exceed the 1∘ reference values.

Orientation (a) (maximum tilting orientated towards the 2 mm lower sheet) uses the geometric effects most efficiently.

The cross tension tests demonstrate a remarkable process robustness.

The joint can be orientated such that the maximum tensile shear strength is aligned with the principal load direction without compromising the axial load-bearing capacity in cross tension.

Due to the extensive use of multi-material systems, future investigations should be extended to multimaterial joints, such as aluminum-steel, where stronger forming effects and a more pronounced geometric anisotropy are expected.
In addition to fully asymmetric tumbling strategies, hybrid process concepts combining an initial asymmetric tumbling phase with a subsequent symmetric tumbling phase could be promising. While the asymmetric phase can create a direction-dependent undercut tailored to the dominant load path, the symmetric step may help to homogenize the joint geometry and reduce local stress peaks. Such a hybrid structure could offer improved load-bearing capacities and should be investigated in future studies.

Acknowledgement

The original version of this paper is available on https://www.scientific.net/MSF.1185.19.pdf

Funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - TRR 285/2 - Project-ID 418701707. The authors thank the German Research Foundation for their organisational and financial support. Data regarding the contents of the publication can be requested at www.trr285.de.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.19.pdf

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H. Q. Ang, "An Overview of Self-piercing Riveting Process with Focus on Joint Failures, Corrosion Issues and Optimisation Techniques," Chin. J. Mech. Eng., vol. 34, no. 1, 2021, doi: 10.1186/s10033-020-00526-3.

D. Li, A. Chrysanthou, I. Patel, and G. Williams, "Self-piercing riveting-a review," Int J Adv Manuf Technol, vol. 92, 5-8, pp. 1777-1824, 2017, doi: 10.1007/s00170-017-0156-x.

M. Merklein, M. Jäckisch, C.-M. Kuball, D. Römisch, S. Wiesenmayer, and S. Wituschek, "Mechanical joining of high-strength multi-material systems - trends and innovations," Mechanics & Industry, vol. 24, p. 16, 2023, doi: 10.1051/meca/2023013.

S. Wituschek, L. Elbel, and M. Lechner, "Versatile self-piercing riveting with a tumbling superimposed punch," Materials Research Proceedings, vol. 28, pp. 1111-1118, 2023, doi: 10.21741/9781644902479-122.

P. Groche, D. Fritsche, E. A. Tekkaya, J. M. Allwood, G. Hirt, and R. Neugebauer, "Incremental Bulk Metal Forming," CIRP Annals, vol. 56, no. 2, pp. 635-656, 2007, doi: 10.1016/j.cirp.2007.10.006.

S. Wituschek and M. Lechner, "Investigation of the influence of the tumbling angle on a tumbling self-piercing riveting process," Proceedings of the Institution of Mechanical Engineers, Part L: Journal of Materials: Design and Applications, vol. 236, no. 6, pp. 1302-1309, 2022, doi: 10.1177/14644207221080068.

S. P. Sundar Singh Sivam, V. G. Uma Sekar, A. Mishra, A. Mondal, and S. Mishra, "Orbital Cold Forming Technology - Combining High Quality Forming with Cost Effectiveness - A Review," Indian Journal of Science and Technology, vol. 9, no. 38, 2016, doi: 10.17485/ijst/2016/v9i38/91426.

M. M. Kasaei, R. Beygi, R. J. C. Carbas, E. A. S. Marques, and L. F. M. Da Silva, "A review on mechanical and metallurgical joining by plastic deformation," Discov Mechanical Engineering, vol. 2, no. 1, 2023, doi: 10.1007/s44245-023-00012-9.

L. P. Troeger and E. A. Starke, "Microstructural and mechanical characterization of a superplastic 6xxx aluminum alloy," Materials Science and Engineering: A, vol. 277, 1-2, pp. 102-113, 2000, doi: 10.1016/S0921-5093(99)00543-2.

DVS/EFB 3480-1 testing of properties of joints - testing of properties of mechanical and hybrid (mechanical/ bonded) joints., Deutscher Verband für Schweißen und verwandte Verfahren e.V. (DVS) / Europäische Forschungsgesellschaft Blech (EFB), Düsseldorf, 2007.

S. Wituschek, "Combination of versatile self-piercing riveting processes," in Materials Research Proceedings, 2023, pp. 125-132, doi: 10.21741/9781644902417-16.

This ePub file was created using machine-generated methods and automated software to convert the source text into an ePub format and may contain formatting inconsistencies from the original version. Formatting may vary depending on the device used.

The original version of this eBook is available on https://www.scientific.net/book/innovative-joining-by-forming/978-3-0364-1990-9

Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 29-39

doi: 10.4028/p-hr6sOv

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2025-12-22

Revised:
2026-02-20

Accepted:
2026-02-23

Online:
2026-04-15

Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

Zambelli Marco 1,a*, Bocchi Sara 1, b, Negozio Marco 2,c, D'Urso Gianluca 1, d, Giardini Claudio 1,e and Antonini Laura 1,f

¹Department of Management, Information and Production Engineering, University of Bergamo, Via Pasubio 7/b, 24044 Dalmine, BG, Italy

a  marco.zambelli@unibg.it, b  sara.bocchi@unibg.it, c  marco.negozio@unipr.it, d  gianluca.d-urso@unibg.it, e  claudio.giardini@unibg.it, f  laura.antonini@unibg.it

Keywords: friction stir welding, L-PBF, additive manufacturing, AA6082, tool geometry.

Abstract

In Friction Stir Welding (FSW) tool geometry plays a critical role in governing heat generation, material flow, and microstructural evolution within the weld. In this study, the feasibility and performance of FSW tools manufactured by Laser Powder Bed Fusion (L-PBF) are experimentally and numerically investigated. A non-conventional FSW tool produced in AISI 316L by L-PBF was designed and compared with a conventional machined steel tool in the welding of AA6082-T6 sheets performed using already optimized process parameters. This was followed by tensile testing and macro- and micro-hardness measurements, and a punctual microstructural analysis. In addition, a 3D thermo-mechanical finite element model was employed to forecast and analyze the temperature distribution, the effective strain and the overall material flow. The results show that the tool manufactured using L-PBF enables FSW joints to achieve mechanical properties and welding efficiency similar to those of the standard tool. Finite Elements Model (FEM) model, in good agreement with experimental results, show that the geometry of the additive tool promotes greater plastic deformation and lower peak temperatures, confirming both the validity of the model and the suitability of L-PBF for the advanced design of FSW tools.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

Among solid-state joining methods, Friction Stir Welding (FSW) is considered one of the most efficient, particularly when applied to aluminum alloys [1] [2]. Unlike fusion based processes, such as TIG or MIG/MAG welding, in which the joint region exceeds the melting temperature and undergoes extensive metallurgical transformation, the FSW process is based on the frictional heat generated by a rotating, non-consumable tool to plasticize the material without never reaching its liquid state. The weld efficiency is generally higher than that of TIG without filler metal, and when compared to TIG with filler or MIG/MAG, even though the joint efficiency is sometimes lower, it offers the advantage that drastically reduces distortion, residual stresses, and typical fusion-related defects such as porosity, hot cracking, and embrittlement within the heat-affected zone (HAZ) [3][4]. Welding aluminum alloys can be challenging, and the growing need to join dissimilar alloys [5], often related to cost-saving reasons, adds further complexity to the technology. Also in this scenario, FSW has proven to be more effective than fusion-based processes, as material positioning and process parameters have a decisive influence on joint quality and help avoid many defects typically associated with TIG welding, in which the latter is strongly affected by filler selection, grain growth in the HAZ, and the presence of porosity and micro-voids. In TIG welding, moreover, variations in welding speed can either improve or deteriorate the mechanical properties depending on the selected values, making the process less predictable [6]. Several comparative studies have shown that, even for AA2xxx and AA7xxx that are typically considered the most difficult to weld aluminum alloys, FSW provides superior mechanical properties in every respect compared to other welding techniques [7] [8]. Not only high-strength alloys benefit from the application of the FSW process, but also the AA6xxx series, that are widely used for their good mechanical properties and corrosion resistance. Nevertheless, these properties, the presence of magnesium and silicon makes them particularly sensitive to hot cracking
when fusion temperature is reached. For this reason, conventional welding processes are often improper. This is the reason FSW has shown itself to be an alternative that can be considered reliable, making it possible to create sound joints with detailed microstructure by the process of dynamic recrystallization. Therefore, the FSW method has shown increasing usage in areas that require highquality welding of aluminum, including the aerospace, automotive, and ship industries. It is essential that a well-designed welding tool be used to create a quality FSW joint, as the tool shape has a direct effect on the material flow and, ultimately, the process stability. The tool is composed of two main elements: the shoulder and the pin, each with its own characteristics and specific tasks. The shoulder remains in contact with the surface of the workpiece, providing most of the frictional heat and confining the plasticized material. Shoulder geometries range from flat and smooth profiles to concave shapes and scroll-type designs, each affecting heat input, downward pressure, and stirring behavior in the upper portion of the weld [9]. The shoulder diameter directly influences the thermal input: larger shoulders generate higher temperatures at the same rotating speed, improving material mixing but also increasing the risk of defects such as excessive flash if the heat becomes excessive.
The pin is responsible for penetrating the material and inducing mechanical stirring within the weld nugget. Its length determines the penetration depth, while its geometry that can be cylindrical, tapered, threaded, or equipped with helical flutes, strongly influences material flow and final microstructural development [10]. Consequently, tool geometry plays a decisive role in determining joint integrity, defect formation, mechanical performance, and microstructural quality.
Complex geometries could enhance the efficiency of the weld joints, concave and grooved shoulders improve material retention and optimized pin designs (threaded, fluted, square), promoting uniform mixing and reducing defects [11]. The fabrication of such geometries is challenging with conventional technologies; however, Additive Manufacturing enables the production of tools with complex shapes. The possibility to create these intricate features and channels improves thermal management, extends tool life, and allows for broad customization without increasing manufacturing complexity [12]. In recent years, increasing attention has been dedicated to the production and maintenance of machining tools through additive manufacturing technologies, particularly Laser Powder Bed Fusion (L-PBF) and Directed Energy Deposition (DED) [13,14]. L-PBF offers high dimensional accuracy, fine microstructures, and the ability to fabricate complex geometries that cannot be produced through conventional machining, especially in the inner part of the tool [15]. These features make L-PBF a promising candidate for manufacturing custom or small-batch FSW tools designed with optimized geometries or internal features.
The production of machining tools, and especially FSW tools, through additive manufacturing has been only partially explored, yet the available results are promising in terms of mechanical performance [16] and the ability to create complex geometries [17].
For these reasons, L-PBF is best suited for rapid prototyping, for the fabrication of special tools, or for experimental contexts where only a limited number of parts with specific functional features are needed. In a traditional industrial setting, however, the high cost of powders and the limited build volume of machines often make the process less economical.
In this context, integrating advanced predictive models becomes essential. Such simulation methods can predict the effects of different tool geometries on the material, reducing prototyping time and cost, and when predictive models reach sufficient accuracy potentially enabling a direct transition to production [18].
For these reasons, the present research involved the design and 3D modeling of a non-conventional FSW tool manufactured using L-PBF. Experimental welds were carried out on AA6082 sheets using both a L-PBF tool and a conventionally turned one, followed by mechanical and microstructural testing. In parallel, welding process simulations were performed using DEFORM-3D software to analyze the heat distribution, material flow, and plastic deformation zones. Finally, the experimental results obtained with the conventional and L-PBF tools were compared with the numerical outcomes to validate the simulation model and assess the compatibility of additive manufactured tools with FSW technology.

The first objective of this study was to determine if and how FSW tools produced by L-PBF are suitable for FSW process of AA6082 and how their performances are compared to conventional tools. A further objective was to assess the extent to which DEFORM-3D simulations can foresee the overall process outcomes, thereby validating the model as a predictive one for advanced FSW application.

Materials and Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

For the experimental campaign, AA6082-T6 sheets were used as the base material. The plates were arranged in a butt configuration and clamped firmly against a backing plate set at a 3∘ tilt with respect to the welding direction [19] to promote proper material flow and a stable joint. The first part of the study relied on a standard C40 steel tool, already used in previous work [20] [21].
Once this reference joint was completed, the same welding procedure was repeated using an additively manufactured tool produced by L-PBF, designed with a specific geometry for the purposes of the study (Figure 1). The L-PBF tool was fabricated using gas-atomized AISI 316L powder as the feedstock, which main properties are summarized in Table 1 Printing took place in an inert atmosphere to limit oxidation, considering the parameters listed in Table 1.

Table 1. Properties of the AISI 316 powder and main printing parameters used for manufacturing the L-PBF tool.

LPB-F PRINTING POWDER PROPERTIES
C | Si | Mn | Cr | Ni | Mo | Fe
<0.03 | <1.00 | <2.00 | Min | Max | Min | Max | Min | Max | Bal.
16.00 | 18.00 | 11.00 | 14.00 | 2.00 | 3.00
PARTICLE SIZE DISTRIBUTION | DRYING
D10 | D50 | D90 | Oven at 70°C for 4 hours (moisture removal)
20.70 μm | 34.70 μm | 60.60 μm
PRINTING CONDITIONS

Layer
thickness
[μm]
Laser power
[W]
Scanning
speed
[mm/s]
Hatch
distance
[μm]
Laser
spot
diameter
[μm]
Build plate
preheating
temperature
[°C]

0.30 | 130 | 900 | 80 | 100 | 80

Layer spreading was performed using an antistatic polymer recoater, and a stripe scanning strategy was adopted. Between successive layers, the scan direction was rotated by 67∘ counterclockwise to reduce thermal stress buildup and improve mechanical response. After printing, the tool was detached from the build plate by wire EDM. No heat treatment was needed before it was used for welding. All the welds, regardless of the considered tool, were performed considering the already optimized process parameters combination for the standard tool, i.e. tool rotational speed ( S ) of 2400 rpm and feed rate (F) of 375 mm/min [22].

[picture]

Fig. 1. L-PBF tool (on the left) and Standard tool (on the right). The shoulder of the L-PBF tool has a concave surface with concentric ribs.

Fig. 1. L-PBF tool (on the left) and Standard tool (on the right). The shoulder of the L-PBF tool has a concave surface with concentric ribs.

For both the welds obtained through the standard and the L-PBF tools, mechanical testing included macro and micro-hardness measurements were performed. Surface hardness was determined using the Rockwell B scale (HRB) following UNI EN ISO 6508. Vickers micro-hardness (HV1000) was measured on cross-sections following a grid layout that satisfies the spacing rules of ASTM EN 9217, both the hardness patterns are shown in Figure 2.

[picture]

Fig. 2. Distribution of the HRB and HV1000 indentations on the specimen.

Fig. 2. Distribution of the HRB and HV1000 indentations on the specimen.

Tensile specimens were prepared in line with the geometry set by UNI 10002. Tests were conducted using a Galdabini tensile test machine, with a 0.2 kN preload applied. Two welding samples were tested for each tool. The crosshead speed was fixed at 1.2 mm/min in accordance with UNI EN ISO 6892-1:2019. A MATLAB script was then used to link the applied force acquired by the tensile machine with the measured strain ( ϵy ) and to produce stress-strain curves.
After mechanical testing, the welded joints were sectioned, mounted in resin, and polished with abrasive papers up to 4000 grit, followed by final polishing with 3μ m and 1μ m alumina suspensions. Microstructural features were examined using a digital optical microscope after a double chemical etching. The first etchant consisted of 25 mLHCl(1 N),25 mLHNO3(65%), approximately 0.2 mL of HF ( 50% ), and 25 mL methanol ( 95% ); afterward, Weck's reagent was applied.
The grain size measurements were carried out following the Heyn method reported in ASTM E11213, using the discretization scheme shown in Figure 3.

[picture]

Fig. 3. Microstructures Acquisition Points Using the Optical Microscope.

Fig. 3. Microstructures Acquisition Points Using the Optical Microscope.

Using a MATLAB code, a microstructure map of the welded zone was reconstructed, displaying the average transversal grain diameter.
The numerical approach adopted in the present work built directly on a thermo-mechanical simulation model, developed in DEFORM-3D, that has already been validated experimentally in a previous study [18]. The reference model is based on a three-dimensional Lagrangian approach, able to track the local evolution of strain, strain rate, and temperature with good spatial resolution, and it adopts an adaptive remeshing strategy to preserve numerical stability and accuracy, as the large plastic deformation inherent to FSW would otherwise lead to excessive mesh distortion. This strategy involves locally regenerating the mesh in the most highly deformed regions. In the simulation, the tool is assumed to behave as a rigid body, while the workpiece is modeled as a visco-plastic material, where plastic flow clearly dominates over elastic effects. The mechanical response of the material is defined using temperature- strain- and strain-rate-dependent flow curves taken from the DEFORM material database.
Heat generation during welding is accounted for through two contributions: friction at the interface between tool and workpiece and heat produced by plastic deformation in the stirred zone. In this way, the model reproduces the thermal cycles that control microstructural evolution. The boundary conditions governing heat transfer by conduction, convection, and radiation were calibrated using experimental data (Table 2), resulting in simulated temperature profiles that deviate from the measured values by less than 10%, as reported in [23].

Table 2. Calibration parameters for DEFORM simulations.

Parameters | Value
Aluminum Emissivity | 0.25
Aluminum Thermal Conductivity [N/(mm·s·°C)] | 450
Friction Coefficient Aluminum-Steel | 0.6
Heat Transfer Coefficient Aluminum-Steel [N/(mm·s·°C)] | 11
Heat Exchange with the Environment [N/(mm·s·°C)] | 0.02
Mechanical Conversion to Heat | 0.8

The thermo-mechanical simulation ultimately provides the complete history of effective strain, strain rate, and temperature at each material point and the same model was used considering fixed process parameters (S2400 rpm and F375) and by changing the geometry of the tool, in order to compare the obtained results also from the simulation point of view.

Results

The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

The welds have a visually satisfying overall appearance, with consistent and clearly defined bead morphology in all conditions. No macro-defects were observed in any of the joints, including tunnel voids, lack of penetration, or root defects.
The graphs reported in Figure 4 show the hardness distribution along the transverse cross-section of the FSW joint, evaluated in terms of both Rockwell B macro-hardness (HRB) and Vickers microhardness (HV1000), as a function of the distance from the weld center.

[picture]

Fig. 4. HRB and HV1000 Hardness Comparison Charts, L-PBF Tool Compared with Traditional Tool at F375 S2400.

Fig. 4. HRB and HV1000 Hardness Comparison Charts, L-PBF Tool Compared with Traditional Tool at F375 S2400.

In both cases, the typical W-shaped profile associated with the FSW process is observed. Hardness shows its maximum in the base material, gradually decreases when approaching the welded zone, and attains a minimum near the transition between the nugget and the TMAZ, before increasing almost symmetrically on the opposite side of the joint. The HRB profile clearly highlights, for both the geometries of the tool, a localized softening region around the center of the joint, with minimum hardness values that vary depending on the tool configuration. The differences observed among the profiles suggest that tool geometry affects less the extent and more the severity of the softening, likely through variations in heat generation and material flow. In particular, The L-PBF tool shows minimum hardness values comparable to those of the conventional tool, however, on average, it exhibits lower hardness in the nugget and TMAZ regions, following a more uniform hardness trend. The HV1000 micro-hardness profile confirms the hardness reduction in the central region of the joint, even if it is evidently more gradual and directly reflects local microstructural modifications, particularly grain size and precipitate condition. The recovery of hardness toward the base material is symmetric, indicating that the observed variations are mainly governed by the thermal field rather than by macroscopic welding defects. In this case, it can be stated that the hardness curves obtained using the conventional tool and the L-PBF tool are practically overlapped.
The tensile test results (Figure 5) obtained using the L-PBF tool show a very consistent mechanical response. The two stress-strain curves exhibit similar trends, with only minor variations in ultimate tensile strength (UTS). This limited UTS variation is a positive indicator of good process repeatability and highlights the stability of the welding conditions, regardless the considered tool geometry.

[picture]

Fig. 5. Tensile Test curves of: L-PBF Tool compared with Standard Tool at S2400 F375.

Fig. 5. Tensile Test curves of: L-PBF Tool compared with Standard Tool at S2400 F375.

When compared with the reference values obtained using the standard Tool, the UTS achieved with the L-PBF tool is fully comparable, indicating that the adoption of an additively manufactured tool does not adversely affect joint strength. The welding efficiency is approximately 68%, with minor variations observed between joints produced using the traditional tool and those welded with the LPBF tool. Overall, the tensile results confirm that the L-PBF tool is capable of producing sound and repeatable joints with mechanical performance comparable to that obtained using the standard one. The FSW process is inherently a thermomechanical process, in which it is important to evaluate the mechanical characteristics of the obtained welds, but it is equally important to analyze how the deformations and heat flow generated by the process are distributed in the material being processed. This is because these characteristics also risk significantly affect the tool wear, reducing the useful life of the tool, leading to pin deformation and thus affecting the quality of the weld. For this reason, the responses of the simulation model in terms of effective strain and temperature were analyzed. The images show the responses of the numerical model in terms of effective strain (Figure 6) and temperature (Figure 7), allowing a direct comparison between the square tool made using L-PBF and the standard tool.

[picture]

Fig. 6. DEFORM simulation results for effective strain: standard tool compared with L-PBF one.

Fig. 6. DEFORM simulation results for effective strain: standard tool compared with L-PBF one.

[picture]

Fig. 7. DEFORM simulation results for temperature: standard tool compared with L-PBF one.

Fig. 7. DEFORM simulation results for temperature: standard tool compared with L-PBF one.

The effective strain maps clearly show that, in the case of the square tool (Figure 6), the weld is characterized by significantly higher levels of plastic deformation, strongly concentrated in the area along the tool-workpiece interface. This behavior indicates greater mechanical stirring action, favored by the non-axially symmetric geometry of the pin, which increases material mixing and the severity of local deformations. In the case of the standard tool (Figure 6), on the contrary, the deformations are lower and more distributed, suggesting a less marked mechanical contribution to the welding process. This difference highlights how the geometry of the tool plays a decisive role in balancing the mechanical and thermal components of the FSW process.
Opposing considerations emerge from the analysis of the temperature maps (Figure 7). In the case of the standard tool (Figure 7), both the tool and the welded material are subjected to more intense thermal cycles, with higher temperatures and more pronounced thermal gradients in the process zone. This indicates a greater incidence of heat generated by friction, which becomes the dominant component of the process.
In the square tool model (Figure 7), on the other hand, the maximum temperatures are lower overall. This suggests that a greater proportion of the energy introduced is dissipated in the form of plastic deformation rather than heat, reducing the extent of the thermal cycles to which both the workpiece and the tool are subjected.
Overall, the images confirm that the square tool tends to promote the mechanical component of FSW, while the standard tool stresses the thermal component of the process. Both approaches have advantages and critical issues: high deformations can improve mixing but accelerate wear or deformation of the pin and more severe thermal cycles can negatively affect the microstructure of the joint and contribute to thermal degradation of the tool.
The same conclusions can be drawn by examining the results obtained experimentally from the measurements of the average grain diameters in the different zones characterizing the FSW joint. From a microstructural point of view (Figure 8), this results into a stir zone characterized by fine, homogeneous grains, which are more extensive in the case of the square tool than in the standard tool.

[picture]

Fig. 8. Increase/Decrease in Grain Size Using the LPBF Tool Compared to the Standard Tool.

Fig. 8. Increase/Decrease in Grain Size Using the LPBF Tool Compared to the Standard Tool.

The average reduction in grain size within the region comprising the nugget and the TMAZ is approximately 33%, with the crystalline microstructure remaining, on average, finer up to about 16 mm from the weld center.
The high mechanical component of the FSW process associated with the square tool, highlighted by the higher effective strain values, promotes intense plastic deformation and dynamic recrystallization phenomena, which are responsible for refining the microstructure and enlarging the area of effectively remixed material. In the case of the standard tool, on the contrary, the greater incidence of the thermal contribution leads to a more limited central zone, in which grain refinement is less extensive, as the process is governed more by friction heating than by severe plastic deformation.

Conclusion

The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

In this work, the feasibility of using additively manufactured Friction Stir Welding tools produced by Laser Powder Bed Fusion was investigated through a combined experimental and numerical approach. Experimental welding trials on AA6082-T6 sheets demonstrated that the L-PBF tool is capable of producing sound joints with no macroscopic defects and with mechanical properties fully comparable to those obtained using a conventional, machined tool. Hardness distribution and tensile test results proved a similar welding efficiency and weld quality between these two tool arrangements.

Numerical simulation results had offered a great benefit in understanding the thermo-mechanical process characteristics and the effect of tool geometry. It had been demonstrated that the tool geometry designed using L-PBF indicated an enhanced degree of plastic deformation with lower maximum temperatures than with the conventional tool; therefore, there was an apparent change in the relative contributions from the mechanical stirring component and the friction component. This difference had resulted in variations in both the predicted grain structure and its homogeneity.
The good agreement between experimental data and simulation results proves the correctness of the simulation model selected for the simulation and the effectiveness of the L-PBF process for FSW tool production. In summary, the obtained results have shown to confirm the effectiveness of the L-PBF process for manufacturing FSW tools and have underlined the potential of this process for future industrial applications. Furthermore, the integration of experimental testing with advanced numerical modeling proves to be a powerful approach for supporting the design, optimization, and future development of innovative FSW tool geometries.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.29.pdf

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 41-53

doi: 10.4028/p-Kld5TA

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-13

Revised:
2026-02-15

Accepted:
2026-02-23

Online:
2026-04-15

Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

Schwichtenberg Nik 1,a*, Tesch Markus 1, b, Sigl Martina E. 1,c, Goeldner Moritz 1, d and Zaeh Michael F. 1,e

1 Institute for Machine Tools and Industrial Management (iwb), TUM School of Engineering and Design, Technical University of Munich, Boltzmannstrasse 15, Garching, Germany

a* nik.schwichtenberg@iwb.tum.de, b  markus.tesch@iwb.tum.de, c  martina.sigl@iwb.tum.de, d  moritz.goeldner@iwb.tum.de, e  michael.zaeh@iwb.tum.de

*corresponding author

Keywords: micro friction stir spot welding, current collectors, busbars, electrical contacting, printed circuit boards.

Abstract

The growing demand for advanced energy storage systems, power electronics, and electrical interconnections requires joining processes that ensure both electrical conductivity and mechanical stability. Applications such as battery cell-internal contacts, flexible busbar connections, and printed circuit board (PCB) conductor interfaces impose strict requirements on material combinations, layer thicknesses, and thermal management, which influence the suitability of the joining processes. Micro friction stir spot welding ( μ FSSW) offers advantages through its low thermal impact as well as its high process robustness and through the bonding of the joining partners in solid state. In the studies presented in this article, transferable simplified specimens were designed to systematically evaluate process parameters for different applications of μ FSSW. The feasibility of a full contact using μFSSW was assessed through electrical resistance measurements, mechanical shear testing, and morphological examinations of the weld zone, as well as axial force measurements during the welding process. The results demonstrate the suitability of μFSSW for copper-based current collector foils in cell-internal contacts, busbar interconnections, and PCB junctions, and they highlight key relationships between the process conditions and the electrical, mechanical, and structural performance. The study also highlights challenges and opportunities for a future industrial implementation.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

Electrical interconnections play a critical role in modern industrial systems, where reliability, low electrical resistance, and mechanical robustness must be ensured across increasingly diverse product portfolios. This is particularly evident in lithium-ion batteries for which performance demands have risen sharply due to the global push towards electrification and energy-dense storage systems. Higher energy densities require thinner current collector foils and more electrode layers, all of which increase the sensitivity of the contacting process to geometric variations [1,2]. And the teardown studies of modern lithium-ion cells, including the Tesla 4680 and the BYD Blade formats, illustrate that the number of electrodes, foil thicknesses, and tab connection quality directly influence electrical losses and thermal behavior [3,4]. These geometric variations necessitate contacting processes with a high robustness towards changing joining setups.
Busbar systems for power electronics and e-mobility have to meet the requirement of growing current levels by adjusting conductor cross-sections or by integrating multilayer copper stacks, which pose additional challenges for joining technologies [5,6]. Laminated busbars and complex multilayer conductor systems necessitate joints with a low electrical resistance and negligible thermal distortion. PCB-based conductor interfaces introduce another challenge, as embedded and surface-mounted copper structures are highly sensitive to local heat input [7].
Established joining processes, such as ultrasonic welding, resistance spot welding, and laser beam welding are widely used but show significant limitations concerning these evolving requirements.

Ultrasonic welding is strongly affected by variations in stack thickness and foil count, which hinder energy coupling and lead to reduced weld uniformity [6]. Laser beam welding offers high precision but introduces steep thermal gradients, risking damage to sensitive battery and PCB substrates and causing inconsistent weld depths in multilayer copper stacks [8-11]. Resistance spot welding suffers from the high thermal conductivity of copper, which makes nugget formation difficult and increases the risk of expulsion or interfacial defects [12].
Micro friction stir spot welding ( μ FSSW) offers a solid-state, low-temperature alternative that directly addresses these challenges of changed joining partner geometries. In this work, the term micro friction stir spot welding refers to the application of friction stir spot welding to ultra-thin sheet and foil stacks with very low thermal mass. The designation "micro" therefore relates to the scale of the thermo-mechanical interaction zone and process conditions rather than to the absolute tool dimensions [13, 14]. As demonstrated by Sigl et al. [1], μ FSSW enables robust joining of thin copper and aluminum current collector stacks while maintaining low contact resistance. Earlier work by Li et al. [15] showed that μ FSSW leads to stable interconnections with limited thermal impact and favorable electrical behavior, making it well-suited for copper-based applications. Recent research emphasizes the remarkable robustness of the process against geometric variations, including changes in foil number, layer thickness, and stack composition. Tool geometry and process parameters influence hook formation, bond width, and static strength, but the process window remains highly tolerant [16-18]. For busbar and multilayer applications, friction-stir-based processes have been shown to achieve copper joints with refined microstructures and a predictable and enhanced conductivity behavior [19]. Given the limitations of conventional joining processes and the increasing geometric complexity of industrial applications, μ FSSW provides a promising alternative for producing stable, reproducible, and electrically efficient joints. However, a systematic comparison across battery current collectors, busbars, and PCB interconnections has not yet been comprehensively addressed. This work therefore investigates the suitability of μ FSSW across these three representative application classes, evaluates its feasibility using simplified application-specific geometries, and analyzes the influence of relevant process parameters. The study is structured into two consecutive experimental investigations.
In the first study, selected application-specific specimens were examined using simplified geometries that reflect the essential joint designs of the respective applications. This feasibility study aimed to achieve full penetration of the multilayer stacks and to identify characteristic effects arising from the different joint configurations and material arrangements.
Building on these results, a second study was conducted using derived analysis specimens that consolidate the characteristic contact features of the investigated applications into a unified geometry. This approach enables a reproducible and statistically meaningful evaluation with a focused assessment of electrical resistance and mechanical strength, providing a consistent basis for analyzing process-property relationships.

Specimen Concept

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

The three selected application cases share one common functional principle. In all configurations, multiple thin conductive elements must be reliably joined within a confined interface to form a lowresistance current collector. However, the geometric characteristics of the joining partners vary significantly across applications. Transferable specimen designs were therefore needed that fulfill both the functional requirements and the differing boundary conditions. For the feasibility study, simplified geometries were derived that represent the essential stack dimensions and joining zones of each application while eliminating construction-specific constraints. These simplified specimens enable a consistent assessment of μFSSW under comparable conditions.
Fig. 1 shows illustrations of the three applications. In lithium-ion battery pouch cells (a), thin copper or aluminum electrode foils are contacted to a metallic arrestor tab to form the internal current collector. Flexible busbars (b) conduct high currents between power-electronic subsystems using laminated copper or aluminum foils enclosed by thin forming sheets, combining a high currentcarrying cross-sectional area with mechanical flexibility. PCB-based conductor interfaces (c) rely on
copper structures on or within FR4 substrates (common substrate material in printed circuit boards), either as surface-mounted foil stacks or as embedded copper baseplates.

[picture]

Fig. 1. Joining geometries of investigated current collector applications:
a) cell-internal contacts, b) flexible busbars, c) PCB collector spots, coated (top) and embedded (buttom).

Fig. 1. Joining geometries of investigated current collector applications: a) cell-internal contacts, b) flexible busbars, c) PCB collector spots, coated (top) and embedded (buttom).

For the feasibility study, only copper-based joints were investigated. This avoided additional effects from dissimilar interfaces and enabled a clear assessment of μ FSSW under controlled, comparable conditions. The selection of copper is justified, as it accounts for a substantial proportion of conductor applications in electrical systems [20]. The specimen configurations derived are presented in Table 1. The materials investigated include OF-Cu R200 and PHC SE-Cu R360, which are widely used in electronic components due to their high electrical conductivity, as well as FR4.

Table 1. Materials and configurations of the application specimens.

Cell-internal contact | PCB collector spot (coated)
Qty | Material | Size in mm |  | Qty | Material | Size in mm
Arrestor | 1 | OF-Cu R200 | 40×20×0.2 | Cover sheet | 1 | OF-Cu R200 | 15×20×0.2
Foil stack | 11 | PHC SE-Cu R360 | 40×20×0.01 | Foil stack | 10 | PHC SE-Cu R360 | 0.01
Bottom sheet | 1 | OF-Cu R200 | 20×20×0.2 | PCB | 1 | FR4 | 30×30×0.8
|  |  |  |  | 1 | ED-Cu | 30×30×0.035
Flexible busbar | PCB collector spot (embedded)
Qty | Material | Size in mm |  | Qty | Material | Size in mm
Forming sheet | 1 | OF-Cu R200 | 40×20×0.2 | Cover sheet | 1 | OF-Cu R200 | 15×20×0.2
Foil stack | 10 | PHC SE-Cu R360 | 40×20×0.01 | Foil stack | 10 | PHC SE-Cu R360 | 40×20×0.01
Forming sheet | 1 | OF-Cu R200 | 40×20×0.2 | PCB | 1 | FR4 | 30×30×0.5
|  |  |  |  | 1 | OF-Cu R200 | ø10×0.5

Based on the simplified geometries used in the feasibility study, a universal analysis geometry was derived to enable reproducible and comparable measurements for electrical and mechanical testing across all application cases. This geometry represents the essential functional elements of the multilayer conductor joints while standardizing overlap width, foil stack composition, and clamping conditions. As shown in Fig. 2, the analysis specimen provides a controlled experimental environment for evaluating mechanical properties such as shear strength as well as electrical characteristics like contact resistance. By eliminating application-specific geometric variability, this specimen allows the influence of process parameters and of the weld morphology on the joint performance to be assessed consistently and with high statistical reliability.

[picture]

Fig. 2. Schematic illustration and dimensions of the analysis geometry.

Fig. 2. Schematic illustration and dimensions of the analysis geometry.

In Table 2, the materials and specimen configurations used in study 2 are summarized. The number of foils in the stack was variably scalable, which allowed for different conductor cross-sections. A bottom sheet was included only for smaller stacks to prevent unintentional welding to the clamping fixture. This adaptable setup ensured consistent testing across a wide range of multilayer configurations.

Table 2. Material and configuration of the analysis specimens.

Qty
Material

Size

in mm

Cover sheet | 1 | OF-Cu R200 | superscript number
Foil stack | 11 | PHC SE-Cu R360 | superscript number
Bottom sheet | 1 | OF-Cu R200 | superscript number

Experimental Setup

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

The welding experiments were conducted in position-controlled mode on a four-axis horizontal milling machine (MCH 250, Gebr. Heller Maschinenfabrik GmbH, Nuertingen, Germany), which had been adapted for friction stir welding. The μ FSSW tool was manufactured from hardened X38CrMoV5-1 tool steel and coated with a TiBZr to improve wear resistance and thermal stability. The detailed tool geometry and dimensions are shown in Fig. 3.

[picture]

Fig. 3. μFSSW tool geometry.

Fig. 3. μ F S S W tool geometry.

For the feasibility study, representative application specimens were welded using the parameter sets listed in Table 3 in total 48 welding specimens. They were chosen based on prior experimental results. The objective of study 1 was to determine parameter combinations that achieve full penetration and stable weld formation for the three application-specific stack configurations. Study 1 was designed
as a feasibility and transferability investigation rather than a parameter optimization. A limited number of representative parameter sets was therefore applied across all application geometries to identify geometry-dependent effects in weld formation.

Table 3. Process parameters for the application trials (study 1).

Parameter Configuration (PC) | Rotational speed n in min⁻¹ | Plunge depth zp in mm | Plunge speed vp in mm/min | Dwell time td in s
PC 1 | 2900 | 0.05 | 20 | 0.25
PC 2 | 2900 | 0.05 | 20 | 0.5
PC 3 | 2900 | 0.75 | 20 | 0.25
PC 4 | 3400 | 0.05 | 20 | 0.25

Based on these results, a study 2 was conducted using the standardized analysis specimen and applying a full factorial experimental design. The variations in the parameters investigated are summarized in Table 4 with in total 384 welding specimens. Each parameter combination was produced twice. The focus was on covering a wide parameter space and identifying general trends rather than determining the variance at individual parameter points. The plunge depth is defined relative to the shoulder surface.

Table 4. Process parameters for the parameter study (study 2).

Rotational speed

n in superscript number

Plunge depth

Zp in mm

Plunge speed

mathematical formula in mathematical formula

Dwell time

mathematical formula in s

2900 | 0.00 | 20 | 0.1
3400 | 0.025 | 30 | 0.15
3900 | 0.05 | 40 | 0.25
4400 | 0.075 |  | 0.5

During all welding trials, the process forces and the moments were recorded using a SPIKE mobile HSK-A63 HD20 110 A sensor (Pro-Micron GmbH, Kaufbeuren, Germany), which was mounted between the spindle and the tool. This system enables real-time acquisition of axial forces as well as bending and torsional moments, providing additional insights into material flow and weld formation. A subset of welded specimens was prepared for metallographic cross-section analysis and examined using a Keyence VHX-7000 digital microscope to evaluate weld morphology, bonding interfaces, and the occurrence of welding defects. Mechanical characterization was conducted as lap-shear testing based on DIN 50154 [21] on a Zwick Roell Z050 testing machine (ZwickRoell GmbH & Co. KG, Ulm, Germany). The electrical performance was evaluated using a four-point resistance measurement setup (see Fig. 2) based on the methodology described by Grabmann et al. [22]. Table 5 lists the applied measurement parameters.

Table 5. Measurement parameters for mechanical and electrical testing.

Tensile testing | Resistance testing
Parameter | Value | Parameter | Value
Test speed vp in mm/min | 2 | Test current I₀ in A | 10
Clamping distance L₀ in mm | 60 | Voltage U₀ in V | 5
Preload F₀ in N | 1 | Ref. temperature tref in °C | 22

This integrated experimental framework enabled a reproducible assessment of the weld formation, the mechanical strength, the electrical resistance, and the process behavior for all investigated μFSSW conditions.

Results of the Application Trials (Study 1)

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

In the application experiments, micro-friction stir spot welds were successfully manufactured for all investigated specimen configurations. Only in one case, namely the PCB collector spot (coated) specimen (PC 4), delamination was observed. This might have occurred due to the high rotational speed. Fig. 4 shows the corresponding cross-sectional images of the welded joints. The welds exhibit characteristic features typical for the μ FSSW process, including the distinct tool imprint and the associated local reduction in specimen thickness, accompanied by material displacement into the surrounding regions. Within the cover sheet, this displacement leads to a localized deformation that forms a void region and induces a hooking effect into the foil stack underneath. These structural characteristics provide essential insights into the material flow and joint formation mechanisms governing the resulting weld quality.

[picture]

Fig. 4. Cross-sections of the application trials of study 1 (PC 3): a) cell-internal contact, b) flexible busbar, c) PCB collector spot (coated) and d) PCB collector spot (embedded).

Fig. 4. Cross-sections of the application trials of study 1 (PC 3): a) cell-internal contact, b) flexible busbar, c) PCB collector spot (coated) and d) PCB collector spot (embedded).

Several irregularities were observed in the weld perimeter and surrounding material zones. Examples are illustrated in Fig 5. These irregularities include the formation of flash around the shoulder contact area, local delamination within the laminated structure, and notches resulting from partial bonding at the outer edge of the weld. While no fractures or breakages of the foils were detected in the examined specimens, such defects cannot be fully ruled out, particularly in regions experiencing pronounced deformation. Likewise, the generation or embedding of micro-debris caused by abrasive interactions during the process may occur, although no distinct particles were identified in the present metallographic sections. These observations highlight the complex interaction between tool penetration, material flow, and layered composite response under the thermomechanical loading conditions characteristic of μFSSW.

[picture]

Fig. 5. Cross-sections of irregularities of the application trials (PC 4):
a) cell-internal contact, b) PCB collector spot (coated) and c) PCB collector spot (embedded).

Fig. 5. Cross-sections of irregularities of the application trials (PC 4): a) cell-internal contact, b) PCB collector spot (coated) and c) PCB collector spot (embedded).

The evaluation of the axial force signals for the different application specimens reveals that all copper-based specimen configurations exhibit a comparable force progression (see Fig. 6) during the μ FSSW process. Such force curves can be clearly associated with the fundamental process phases, namely the plunge phase during which the tool penetrates and plasticizes the material, the dwell phase during which frictional heating and material consolidation dominate, and the retract phase in which the tool is withdrawn and the joint solidifies, as presented in the figure. For the investigated applications cell-internal contact, flexible busbar, and PCB collector spot (embedded), the extracted curves display nearly identical trends, indicating a high level of repeatability and demonstrating that the axial force response is predominantly influenced by the underlying thermomechanical process rather than by moderate geometric or thickness variations between the specimen designs. In contrast, the PCB collector spot (coated) application exhibits a noticeably increased maximum axial force. This behavior can be attributed to the reduced deformation capability of the substrate plate.
Beyond the overall curve progression, characteristic points relevant for process interpretation can be identified, including the initial force rise during the plunge phase, the force stabilization and material consolidation during the dwell phase, and the distinct drop associated with the retract phase. Only the force discontinuity that typically occurs when the tool shoulder reaches the material, often visible as a sharp force jump, was not always clearly pronounced. This variation may result from local surface conditions or slight differences in contact mechanics; however, it does not impede the overall comparability of the curves.
A deviation from this consistent behavior is observed for the PCB collector spot-coated application. In contrast to the purely copper-based specimens, this configuration incorporates an FR4 substrate, whose mechanical stiffness, thermal conductivity, and deformation behavior differ considerably from those of metallic substrates. These material differences may dampen or distort segments of the axial force curve, resulting in a reduced transferability of the characteristic force progression. This effect was considered during the design of the analysis geometry to ensure that material-specific deviations do not compromise the identification of general process features.

[picture]

Fig. 6. Axial force curves of the application trials (PC 3):
1: plunge phase, 2: dwell phase, 3: retract phase.

Fig. 6. Axial force curves of the application trials (PC 3): 1: plunge phase, 2: dwell phase, 3: retract phase.

Results of the Parameter Study (Study 2)

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

To evaluate the influence of the main process parameters on the resulting quality characteristics, a correlation analysis was conducted using both the Pearson correlation coefficient and the Spearman rank correlation. The application of these two methods enables the detection of different types of statistical relationships. The Pearson correlation coefficient describes linear dependencies between metric variables and assumes an approximately linear relationship or normally distributed data [23]. Since thermo-mechanical joining processes frequently exhibit parameter regions with near-linear scaling behavior, Pearson is well suited to identify such relationships. In contrast, the Spearman correlation is based on ranked data and captures monotonic, but not necessarily linear relationships, making it more robust against outliers and able to represent nonlinear trends [24]. This is particularly advantageous because joining processes often exhibit saturation effects or threshold behaviors. The combined use of both approaches therefore provides a comprehensive view of the underlying parameter-response structure (see Table 6).

Table 6. Correlation matrix according to Pearson and Spearman for the process parameters and the electrical and mechanical properties.

| Pearson correlation/ p-value | Spearman correlation/ p-value
Process parameter | El. resistance | Tensile force | El. resistance | Tensile force
Plunge depth | -0.68/ 5.1×10⁻⁵³ | 0.57/ 1.6×10⁻³³ | -0.69/ 5.4×10⁻⁵⁶ | 0.58/ 7.9×10⁻³⁵
Rotational speed | -0.31/ 1.3×10⁻⁹ | 0.37/ 1.2×10⁻¹³ | -0.30/ 3.0×10⁻⁹ | 0.38/ 6.3×10⁻¹⁴
Dwell time | -0.22/ 1.7×10⁻⁵ | 0.10/ 0.065 | -0.23/ 2.3×10⁻⁵ | 0.13/ 0.023
Plunge speed | 0.07/ 0.13 | -0.18/ 3.9×10⁻⁴ | 0.07/ 0.13 | -0.18/ 5.1×10⁻⁴

Both correlation methods reveal consistent trends, confirming that plunge depth and rotational speed are the primary factors influencing electrical resistance and tensile strength. Plunge depth exhibits the highest absolute correlation coefficients, showing a negative relationship with electrical resistance and a positive relationship with tensile strength. This is consistent with the physical expectation that deeper tool engagement promotes enhanced plastic deformation and improved bonding conditions, thereby facilitating effective metallic contact formation.
Rotational speed shows moderate but stable correlations with both target metrics, reflecting its role in frictional heat generation, material plasticization, and joint consolidation. From a statistical perspective, the associated correlation coefficients are accompanied by low p -values, indicating that these relationships are unlikely to arise by chance. However, statistical significance is interpreted in conjunction with effect size and physical plausibility rather than as a binary criterion.
The dwell time displays lower correlation coefficients than the plunge depth but reaches values comparable to those of the rotational speed for the tensile strength. Although the corresponding pvalues indicate only marginal statistical significance, this suggests that the dwell time contributes measurably to a process refinement rather than acting as a dominant control parameter within the investigated process window.
In contrast, the plunge speed shows only very weak correlations with the electrical resistance. While these correlations may reach statistical significance due to the relatively large sample size, their low effect magnitude indicates limited practical relevance. Consequently, the influence of the plunge speed on the electrical performance is considered secondary.
Based on these findings, the subsequent graphical evaluation focuses on plunge depth and rotational speed as the primary variables governing the joining outcome. These parameters exhibit the strongest and most physically meaningful relationships with the quality metrics and therefore provide a robust basis for analyzing process-property interactions. This rationale underpins the decision to focus further investigations on these two factors.
The boxplot representation (median, upper and lower quantile) of the electrical resistance and mechanical tensile strength for the selected process parameter combinations provides a clear visual indication on how the plunge depth and the rotational speed influence the joining performance (see Fig. 7). As observed, increasing both plunge depth and rotational speed leads to a systematic improvement in joint quality, reflected by a decrease in electrical resistance and an increase in tensile strength. These trends are consistent with the outcomes of the correlation analysis, which identified both parameters as the primary drivers of the process behavior due to their strong monotonic and linear relationships with the target metrics.

[picture]

Fig. 7. Boxplots of electrical resistance (top) and tensile strength (bottom) for different combinations of plunge depth and rotational speed.

Fig. 7. Boxplots of electrical resistance (top) and tensile strength (bottom) for different combinations of plunge depth and rotational speed.

A notable feature in the graphical evaluation is the distinctly larger spread of values associated with the process parameter combination of a plunge depth of 0.025 mm and a rotational speed of 3400 min−1. For this specification, a noticeable increase in scatter is observed in both electrical resistance and mechanical tensile strength. The evaluation across the remaining process parameters, dwell time and plunge speed, reveals that certain parameter settings, especially those involving higher plunge speeds and longer dwell times, lead to substantial fluctuations. In these cases, electrical resistance can vary by more than 100 mΩ, while tensile strength differences exceed 100 N . This pronounced variance indicates that the 0.025 mm/3400 min−1 combination is particularly sensitive to changes in the additional parameters and results in a higher process variability in the joining performance.
In contrast, the parameter sets associated with larger plunge depths and higher rotational speeds show significantly narrower distributions in both resistance and tensile strength. This reduced variance suggests more stable thermomechanical conditions, improved material consolidation, and overall more reliable process outcomes.
Overall, the combined interpretation of the boxplots and the correlation results highlight the central role of plunge depth and rotational speed in determining both the absolute quality values and the
consistency of the joining process. These findings confirm the suitability of focusing subsequent investigations on these two dominant parameters.

Discussion of study 1: Application-specific feasibility and joint robustness

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

The results of the first study demonstrate the suitability and robustness of μ FSSW across the investigated application-specific joint designs. Full penetration of the multilayer stacks was achieved for all current collector applications using an identical tool geometry and a common set of process parameters, indicating a resilience with respect to joint design and stack configuration.
Metallographic cross-sections reveal characteristic features of μFSSW, including tool imprint formation, local thinning, and material displacement around the weld zone, which are consistently observed across all applications. These findings indicate stable material flow conditions and underline the suitability of μFSSW for heterogeneous copper-based current collector designs.
A delamination was observed in the coated PCB collector configuration. However, it cannot be conclusively determined whether this effect originated during welding or occurred during specimen preparation. While this observation suggests a potential sensitivity of FR4-based substrates, it does not allow for an unambiguous attribution to the joining process itself.

Discussion of study 2: Parameter influence and process-property relationships

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

The second study confirms that plunge depth and rotational speed are the dominant process parameters governing the joint quality in μFSSW, consistent with established thermomechanical joining principles [1,2]. Increased plunge depth enhances material engagement and bonding, while higher rotational speed promotes frictional heating and plasticization, resulting in reduced electrical resistance and increased tensile strength.
Dwell time and plunge speed exhibit lower correlation coefficients and therefore act mainly as tuning parameters within the investigated process window. Although their influence is statistically detectable, their effect magnitude remains limited compared to the dominant parameters. Axial force analysis further corroborates these findings. Comparable force-time signatures for copper-based specimens indicate stable process mechanics.

Conclusion and Outlook

The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

This study demonstrates that μ FSSW enables the reliable formation of fully consolidated joints across various copper-based current collectors, with plunge depth and rotational speed identified as the dominant parameters influencing electrical and mechanical performance. The combination of correlation analysis, metallographic examination, and axial force evaluation reveals consistent process behavior among the copper-based applications, while deviations observed in the coated PCB configuration highlight the role of substrate-specific thermo-mechanical properties.
Future work should focus on a more detailed assessment of deformation phenomena, including the extent of thinning, surface indentation and subsurface material displacement, in order to better understand their relevance for electrical and mechanical functionality. Furthermore, investigating internal irregularities, such as voids or delamination fronts, will be essential for establishing robust quality criteria. In addition, the influences of varying foil thicknesses and stack heights are to be investigated.
Acknowledgements. This work was funded by the German Federal Ministry of Research, Technology, and Space (BMBFTR) under grant 03XP0545. We would like to thank the BMBFTR and the project management organization Julich (PTJ) for their support. We thank Jana Hauptmann for the sample preparation und Eren Günal for assisting during the welding experiments.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.41.pdf

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 55-68

doi: 10.4028/p-IFcFC6

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-16

Revised:
2026-02-16

Accepted:
2026-02-24

Online:
2026-04-15

Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

Aron Ringel 1,a*, Ahmet Zahid Demirezen 1, b, Julika Hoyer 2,c, David Bailly 1, d, Simon Kammerloher 2,e, Wolfram Volk 2,f and Junhe Lian 1, g

1 Institute of Metal Forming (IBF), RWTH Aachen University, Intzestraße 10, 52072 Aachen, Germany

2 Chair of Metal Forming and Casting (utg), Technical University of Munich, Walther-Meißner-Straße 4, 85747 Garching, Germany

a* aron.ringel@ibf.rwth-aachen.de, b  ahmet.zahid.demirezen@ibf.rwth-aachen.de, cjulika.hoyer@tum.de, d david.bailly@ibf.rwth-aachen.de, e simon.kammerloher@tum.de, f wolfram.volk@utg.de, gjunhe.lian@ibf.rwth-aachen.de

Keywords: continuous compound casting, multi-material, rolling, FE-simulation.

Abstract

Multi-material components that consist of copper and aluminum enable the combination of advantageous mechanical, thermal, and electrical properties at competitive cost. While roll bonding is an efficient-solid state joining technique, its implementation requires fully processed, cold-rolled strip material from two process routes. Continuous compound casting in contrast offers a more efficient approach by combining aluminum and copper during casting, followed by flat rolling in a single process route. However, the differences in flow stress between the metals can cause nonuniform elongation and therefore significant shear stresses at the interface during rolling. These stresses may lead to a delamination of the compound if process conditions are not well controlled. This study investigates whether a geometrically structured interface, introduced during compound casting, can contribute to withstanding interfacial shear stresses through mechanical interlocking. In finite element simulations varying process parameters including height reduction, initial temperature, rolling speed ratio, and pass schedule were examined. Results show that a structured interface can effectively resist shear stresses at the copper-aluminum boundary, thereby improving joint stability during deformation. Furthermore, the strain distribution as well as the fluctuation of the shear stresses can be controlled by the process parameters. The findings indicate that the mechanical interlocking by a geometric interface combined with optimized process parameters can enhance the rolling of compound-cast copper-aluminum composites.

Introduction

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The integration of different material classes in multi-material components allows for the advantageous combination of their respective mechanical properties. Joining copper ( Cu ) and aluminum ( Al ) results in an attractive combination of high mechanical strength, corrosion resistance, and favorable thermal and electrical conductivity, all at a competitive cost [1]. However, joining Cu and Al using classic welding techniques, i.e., fusion welding, is problematic due to significant differences in their physical and chemical properties. Furthermore, the formation of brittle intermetallic compounds (IMC) can significantly reduce joint strength. A common alternative is explosive welding, which is typically employed for batch production and introduces safety risks due to the need for individual preparation and handling of explosives for each weld [1].

In contrary, roll bonding offers a high-throughput, continuous manufacturing while avoiding these safety hazards. The joint formation mechanism can be described by the film theory: During deformation, the contaminant layers present on the surfaces fracture. As a result, the base material extrudes through these cracks and coalesces with the opposing base material [2]. This mechanism implies that metallurgical bonding occurs only after surpassing a threshold deformation level. Beyond this threshold, the bond strength increases with further deformation until reaching the strength of one base material [3]. Reported threshold height reductions in roll bonding vary depending on the
materials and the process parameters: values range from 40% [4], over 60% [5], up to 80% [6] height reduction during the decisive first rolling pass. Larger initial thicknesses require higher threshold deformations or yield lower bond strengths [7].

To ensure effective bonding, it is essential to remove the contaminant layers by thorough surface preparation. This is typically achieved through chemical cleaning processes such as washing and degreasing [8]. In addition, scratch brushing is employed to eliminate oxide layers and induce work hardening, which promotes surface cracking and thus enhances the coalescence of the base materials [9]. An increased surface roughness has also been shown to improve bond quality [7].

While roll bonding itself is a pressure welding process, the introduction of heat triggers diffusion processes, which can improve the bonding by forming IMC at the Cu/Al interface [6]. Therefore, cold roll bonding followed by an additional annealing is common to improve the interface strength [8,10,11,5]. Other studies use elevated temperatures during rolling itself [6,12]. Regardless of method, the IMC formation at the Cu/Al interface is inevitable [6]. As the IMC evolution depends on diffusion mechanisms, the temperature and time during heat treatment and rolling is decisively affecting the bonding [13]. Post-rolling annealing can initially enhance interfacial diffusion and bond strength but may lead to deterioration if unfavorable IMCs form thereafter [5]. In general, lower temperatures and short heat treatment times favor stronger bonds between Cu and Al layers [10]. Above 250∘C, the growth of IMC accelerates distinctly [8]. Several scientific studies identify 300∘C as an optimal temperature for achieving an uniform distribution of deformation with minimal formation of brittle IMC [5,11]. Conversely, processing at temperatures between 400−500∘C leads to severe development of detrimental IMC [5,11].

Another challenge arises from significantly higher flow stresses of Cu compared to Al during deformation, resulting in unequal thickness reductions for each layer [14]. This can be mitigated either by adjusting initial layer heights or employing asynchronous rolling where roller speeds differ. In [9] the Cu layer was in contact with the faster roller, while in [13], the Cu touched the slower roller. Surprisingly, both found, that a ratio of 1.3 was improving the interface bonding. Furthermore, a higher rolling speed tends to reduce the bond strength [5,7].

Despite its efficiency compared to other methods, roll bonding still necessitates multiple processing steps before achieving a final composite product. This includes separate casting and rolling processes and the extensive surface preparation. To streamline this process chain while reducing energy consumption, direct composite casting approaches such as hybrid casting have been explored. One option is twin-roll casting, where a solid Cu strip passes through an Al melt pool so that Al solidifies onto its surface. This technique produces thin and uniform diffusion layers due to brief exposure at high temperature together with large interface deformation during subsequent rolling [15]. Additional rolling passes combined with annealing further influence the bond strength [16], with promising results reported at 300∘C and 40% height reduction [17].

Advancing this concept further involves continuous compound casting, in which both components are joined directly during casting (Figure 1 a)). First, the Cu melt solidifies in an open graphite mold with a copper cooler using a horizontal continuous casting process. The solidified rod is pulled through the mold by rollers. Then, Al melt is poured onto the solid Cu rod and solidifies in the second part of the mold. During the solidification, an initial metallurgical bonding between Al and Cu can be achieved for the subsequent rolling process. Since no contaminant or oxide formation can get onto the surfaces, common cleaning steps like degreasing or brushing become unnecessary [18]. Plates from a horizontal continuous compound casting process showed best shear strength when rolled at 300∘C with a first-pass height reduction of 55% [18]. Insufficient height reduction led to delamination during rolling. It is assumed that during rolling, the shear stresses can exceed the initial joint strength created by compound casting. When there are not enough freshly coalesced and bonded surfaces formed to compensate, the part may delaminate.

[picture]

Fig. 1. Continuous compound casting a) without and b) with geometrically designed interface.

Fig. 1. Continuous compound casting a) without and b) with geometrically designed interface.

The occurring shear stresses can be influenced by factors including the flow stress difference, and thus the temperature and strain rate, the height reduction, the speed ratio between the rollers, the friction coefficient, and the layer thickness. A previous study has explored a geometric design of the interface to counteract the adverse shear stresses [19]. The interface structure is introduced during compound casting by a structured graphite mold (Figure 1 b)). This allows to use both, a metallurgical bond together with a mechanical interlock.

Given that optimum roll bonding conditions for metallurgical bonding between Al and Cu are well established, this study aims to investigate how varying process parameters affect mechanical interlocking between layers. The hypothesis is, that the combination of metallurgical and mechanical bonding would either enhance the overall bond strength or reduce required threshold deformation.

Material and Methods

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Materials.

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

In this study, the Cu(Cu−ETP) and Al (A199.7) material was cast at utg, TUM, Munich. The flow curves for the materials were measured by compression tests at IBF (Figure 2). The range of the resulting flow stresses of Al is significantly smaller than that for Cu . Thus, no flow stress levelling by choosing a certain temperature seems possible. Therefore, different height reduction and thus elongation of the layers and an inhomogeneous strain distribution between the compound materials is expected.

[picture]

Fig. 2. Flow curves measured in compression test at IBF for a) Al and b) Cu at different temperatures and strain rates.

Fig. 2. Flow curves measured in compression test at IBF for a) Al and b) Cu at different temperatures and strain rates.

FE-Rolling Simulation.

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The rolling process was simulated using Abaqus 2017 with a dynamic explicit solver, incorporating displacement-temperature coupling to represent the thermo-mechanical interactions during deformation. A two-dimensional plane strain assumption was used, as width spread effects were considered negligible for the investigated geometry. The assembly utilized in the simulation is depicted in Figure 3.

The rolling material consisted of Al and Cu plates, with overall dimensions of 120 mm in length and 24 mm in thickness. At the interface, a quadratic surface structure was designed featuring ribs and channels. Each square segment had a side length of 4 mm , resulting in a height of 10 mm in the channels and 14 mm at the ribs for each material. A fillet radius of 0.5 mm was applied to all corners to reflect realistic manufacturing conditions. To ensure that the rolls bite the part, particularly at large height reductions, the front edge of the slab was machined with a chamfer measuring 20 mm in length and tapering from nominal thickness to 8 mm at its thin end. The material properties for Al and Cu layers were assigned based on the flow stress data obtained from compression tests (Figure 2). The Poisson's ratios were set to 0.33 for Al and 0.343 for Cu . An inelastic heat fraction of 0.9 was used for both materials to account for the conversion from plastic work into heat during deformation. Additional temperature-dependent material parameters, i.e., thermal conductivity, density, elastic modulus, coefficient of thermal expansion, and specific heat capacity, were sourced from Springer Materials for Al,Cu, and the work roll material X40CrMoV5 [20-22]. The work rolls themselves featured an outer diameter of 410 mm and were modeled as rigid bodies that remained thermally active throughout the simulation. Guidance along the rolling direction was provided by a roller table positioned with an offset of 15 mm from the highest point of the lower roller surface.

[picture]

Fig. 3. Setup of the rolling simulation.

Fig. 3. Setup of the rolling simulation.

The movement along the rolling direction was achieved via a pusher in combination with a springdashpot system. The pusher was needed to overcome the bite condition in simulations with large height reductions despite the chamfer on the edge. The spring dash-pot system was used to handle the rolling material in pass schedules with several passes by reverse rolling (spring stiffness: 1E-005); dashpot coefficient: 0.01 ). The process started with a 5 s waiting time representing the transfer time from the furnace to the rolling mill. By this, the initial temperature gradients could develop realistically before rolling with a rolling speed of 3.0 m/min. If several rolling passes were needed, the rolling direction was reversed with an interpass time of 5 s . After the rolling, a 5 s time for air cooling was assumed.

Displacement-temperature interactions between all components were fully modeled. Free surfaces exchanged heat with ambient air by convection using a film coefficient of 10 W/(m2 K) (equivalent to 0.01t/(s3 K) ) at a sink temperature of 20∘C. Radiative heat loss was included via emissivity values set at 0.11 for Al and 0.6 for Cu . Thermal contact conductance between the rolling material and both rollers or table surfaces was modeled using tabular data: The maximum conductance values reached
4,000 W/(m2 K) at zero clearance but decreased linearly to zero at clearances beyond 0.05 mm . Initial temperatures for both table and rollers were set uniformly at 20∘C. A frictionless contact was assumed between the rolling material and the roller table. For all other friction-based contacts, the Coulomb friction model was employed. Between the rolling material and rollers, a friction coefficient of μ=0.25 was applied. No metallurgical bond strength was assumed between the Cu and Al layers. However, to account for the initial bonding from the compound casting process, relatively high friction values were selected: An interfacial friction coefficient of μ=0.5 between Cu and Al , and a value of μ=1.0 for self-contact within either metal. The thermal interaction across the Cu−Al interface was modelled with a maximum conductance value of 4,500 W/(m2 K) and the same clearance values as described above.

To assess the influence of the friction coefficient between the rollers and the rolling material, it was varied between 0.2 to 0.4 , which is the typical range for hot or warm rolling. For μ=0.2, unrealistically high pushing forces were needed to overcome the bite condition. With up to μ=0.4, the average strain did increase only up to 4.5% for Al and 7.5% for Cu . The fluctuation of the shear stresses in the interface between Al and Cu increased by only 2.65%. Therefore, μ=0.25 seems to be a valid assumption. Furthermore, the interfacial friction coefficient between Al and Cu was reduced down to μ=0.25. While the average strain did only decrease by 0.4% for Al and 1.7% for Cu , a significant decrease of the interfacial shear stress of 31.5% was calculated. Thus, the interfacial shear stress is strongly depending on the interfacial friction.

The finite element discretization employed coupled temperature-displacement plane strain elements (CPE4RT) for the compound material and the rollers. For the rollers an approx. element size of 3.75 mm was chosen, resulting in 3410 elements per roll. A sensitivity analysis was performed to select an element size for the compound material that balanced adequate accuracy against computational cost. At the edge in contact with the roller, a mesh size of 1.0 mm was selected. At the interface, an element size of 0.25 mm was chosen. A biasing scheme was used over the height. This resulted in 10429 elements for Al and in 10662 elements for Cu . To optimize computational efficiency, semi-automatic mass scaling was applied throughout all simulations with a target increment time of 1E-006 s. Additional tests using increment times of 1E-005 s and 1E-007 s revealed that while a 1E-007 s produced comparable results at the expense of increased computation time, 1E-005 s resulted in excessive deformation artifacts.

Results and Discussion

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Influence of the Surface Structure.

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

The effect of the geometric interface is shown in comparison to a simulation with a flat interface in Figure 4 for 50% height reduction and 300∘C initial temperature. As there is no metallurgical bond strength modelled, the compound material with the flat interface separated (Figure 4 a)). Furthermore, the Al layer is approx. 3 mm thinner and therefore elongated significantly more. With the geometric interface, the layers stuck together during and after rolling and therefore elongated similarly. The average difference in thickness is only approx. 1.2 mm . Due to the different final length of the layers, an upward curvature of the rolling material is observed.

[picture]

Fig. 4. Compound rolling with a) flat and b) geometric interface at 50% height reduction and 300∘C initial temperature.

Fig. 4. Compound rolling with a) flat and b) geometric interface at 50 % height reduction and 300 ∘ C initial temperature.

To assess the influence of the geometry on the stresses at the interface, the frictional shear stresses (CSHEAR) along with the contact pressure (CPRESS) were evaluated for every node of the interface and every 0.2 s of the rolling process. The data was then matched with the position along the rolling direction and plotted along the roll gap as in Figure 5. A filtering by averaging the values within a 1 mm distance was performed to receive the darker trend lines in the diagrams. Figure 5 a) shows the results for the case without the geometric interface. As expected, the stresses start from the roll gap entry and stop at the roll gap exit. As rolling is a compressive deformation process the contact pressure is positive. The shear stresses of the averaged trend line values increase in the beginning until a maximum of 23.1 MPa and drop to zero at the neutral point. The negative section with a minimum of -20.4 MPa appears, before the shear stresses get back to zero after the roll gap exit. This clear trend is not observable for the geometric interface. Here, the shear stresses strongly fluctuate throughout the roll gap and extend even beyond the contact length. The averaged values of the trend line show a maximum of 3.8 MPa and a minimum of -12.0 MPa .

[picture]

Fig. 5. Comparison of contact pressure and frictional shear stress for a) flat interface b) geometric interface along the rolling direction.

Fig. 5. Comparison of contact pressure and frictional shear stress for a) flat interface b) geometric interface along the rolling direction.

To assess the fluctuation of the original shear stress data, fshear  was calculated based on all n data points of the shear stress σshear  in the interface as in Eq. 1.

fshear =1n∑(σshear ,n)2.(1)

Thus fshear  considers the shear stress of all interface nodes of both layers within an approx. distance of 0.25 mm (corresponding to mesh size) and over all time frames sampled in 0.2 s intervals. In simulations, the fluctuation of shear stresses was 11.1 MPa for the flat interface, whereas it increased to 17.24 MPa with the structured interface. This indicates that higher peaks of shear stress occur due to the geometric interface structure. However, despite these elevated stresses, no separation of the layers was observed. The interface structure successfully endured the shear stresses and maintained effective mechanical interlocking.

These results suggest that the geometric interface compels both layers to undergo more similar thickness reduction and elongation during rolling. At the same time, it increases shear stresses at the interface because slippage between layers is prevented. If a metallurgical bond would be added to the geometric interface, lower shear stresses may be advantageous for bond strength. Therefore, further investigation into optimized process parameters is necessary to reduce interfacial shear stress.

Influence of Height Reduction and Temperature.

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

A parameter study with varying height reduction εh={20,30,40,50,60}% and initial temperature Ti={250,300,350,400,450}∘C was performed. Figure 6 shows five example final geometries for the different initial temperatures at 50% height reduction. Here, the strains seem to increase with rising temperature, especially for Al.

[picture]

Fig. 6. Final shape with strain distribution after compound rolling at a height reduction of 50% for a) 250∘C, b) 300∘C, c) 350∘C, d) 400∘C and e) 450∘C initial temperature.

Fig. 6. Final shape with strain distribution after compound rolling at a height reduction of 50 % for a) 250 ∘ C , b) 300 ∘ C , c) 350 ∘ C , d) 400 ∘ C and e) 450 ∘ C initial temperature.

The strain data of each element was extracted from the final shape of the compound material for all simulations. The average strain is displayed in Figure 7 a). For the Al part, the average strain shows a minimum at 300∘C. For the Cu part, a maximum average strain can be observed at 300∘C, even though the change is significantly smaller compared to Al. Thus, the strain difference between Cu and Al is smallest at 300∘C, which corresponds to a more homogeneous deformation of the two layers.

[picture]

Fig. 7. Influence of the temperature and height reduction on a) average strain and b) fluctuation of shear stress of all elements of the rolled compound material.

Fig. 7. Influence of the temperature and height reduction on a) average strain and b) fluctuation of shear stress of all elements of the rolled compound material.

In contrast to that, higher temperatures and smaller height reductions reduce the fluctuation of the shear stresses in the interface between Al and Cu and would therefore be beneficial to reduce stresses on a metallurgical bond (Figure 7 b)). However, the geometric interface is effectively avoiding any significant delamination despite the comparably high shear stresses at a rolling temperature of 300∘C. This gets evident by the resulting geometry of the compound material after rolling with different height reduction at 300∘C as it can be seen in Figure 8.

[picture]

Fig. 8. Final shape with strain distribution after compound rolling at an initial temperature of 300∘C and a) 20%, b) 30%, c) 40%, d) 50% and e) 60% height reduction.

Fig. 8. Final shape with strain distribution after compound rolling at an initial temperature of 300 ∘ C and a) 20 % , b) 30 % , c) 40 % , d) 50 % and e) 60 % height reduction.

Influence of Asynchronous Rolling.

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The effect of asynchronous rolling on the compound material was analyzed by scaling the angular velocity α˙ of the upper roller by the factor of the speed ratio given in Eq. 2.

rα˙=α˙upper α˙lower .(2)

For the upper roller being 1.3 times faster than the lower roller, rα˙=1.3, while for the lower roller being 1.3 times faster than the upper roller, rα˙=0.77. The resulting geometry of the simulations with the strain distribution is shown in Figure 9.

[picture]

Fig. 9. Influence of the rolling speed ratio of a) 0.77,b)1.0, and c) 1.3 on the strain distribution of the rolled compound material.

Fig. 9. Influence of the rolling speed ratio of a) 0.77 , b ) 1.0 , and c) 1.3 on the strain distribution of the rolled compound material.

From the simulations, the strains and shear stresses were extracted. Figure 10 a) shows the resulting average strain with the standard deviation given as error bars for Al and Cu .

[picture]

Fig. 10. Influence of the asynchronous rolling speed ratio on the a) strain distribution and b) on the fluctuation of shear stress of all elements of the rolled compound material.

Fig. 10. Influence of the asynchronous rolling speed ratio on the a) strain distribution and b) on the fluctuation of shear stress of all elements of the rolled compound material.

A higher speed ratio reduces the average elongation of Al , so that the stages between Al and Cu increasingly level out. At a ratio of 1.3, the difference between strains of Al and Cu is smallest. In contrast to that, the fluctuation of shear stress at the interface seems to be smallest for the ratio of 0.77 (Figure 10 b)). However, no clear trend is observable, as a maximum was found for a ratio of 1.0 .

Additionally, it was investigated, whether the rolling speed ratio was a tool to influence the resulting curvature of the compound material. Therefore, the geometry of the Al and Cu surfaces were interpolated and the curvature calculated. No significant effect of the speed ratio on the curvature was overserved as shown in Figure 11. The curvature remained nearly constant, with values of 0.00441/mm for Al and 0.00461/mm for Cu .

[picture]

Fig. 11. Influence of the rolling speed ratio on the curvature of the rolled compound material.

Fig. 11. Influence of the rolling speed ratio on the curvature of the rolled compound material.

Influence of Pass Schedule.

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To assess the influence of the pass schedule, the total height reduction of 50% was separated into two and three passes as shown in Table 1.

Table 1. Heights and height reductions for used pass schedules.

Pass schedule | Passes | h₁ [mm] | εh,1 [%] | h₂ [mm] | εh,2 [%] | h₃ [mm] | εh,3 [%]
A | 1 | 12.0 | 50.0 |  |  |  |
B | 2 | 16.8 | 30.0 | 12.0 | 28.6 |  |
C | 3 | 19.2 | 20.0 | 15.3 | 20.3 | 12.0 | 21.06

In the simulations, the passes were modelled as reversing rolling without reheating and an interpass time of 5 s . The resulting geometry with the occurring strains is displayed in Figure 12 a ) - c). With increasing number of passes, the elongation of the lower Al layer is increasing, so that a free Al end is showing (red circle).

[picture]

Fig. 12. Influence of pass number for a) 1 , b) 2 , and c) 3 passes on the geometry and strain distribution of the rolled compound material; d) average strain for schedules A to C .

Fig. 12. Influence of pass number for a) 1 , b) 2 , and c) 3 passes on the geometry and strain distribution of the rolled compound material; d) average strain for schedules A to C .

This is also evident from the average strain (Figure 12 d )). With increasing number of passes, the average strain for Al is increasing, while it stays constant for Cu . This might be due to the fact, that the part cools more with more passes, as the process time increased, which leads to a higher difference in flow stress and thus more dissimilar elongation of the materials (Figure 13 a)). While for pass schedule A, the final temperature after rolling is 270∘C, it is only 223∘C for pass schedule C . The fluctuation of the shear stress is slightly higher, when using only one pass (Figure 13 b)). The pass schedule with three passes (C) appear to show higher shear stresses in the third pass compared to the schedule with two passes.

[picture]

Fig. 13. Influence of number of passes on the a) average temperature and b) fluctuation of shear stress of all elements of the rolled compound material.

Fig. 13. Influence of number of passes on the a) average temperature and b) fluctuation of shear stress of all elements of the rolled compound material.

Preliminary Experiments for Future Validation

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Continuous compound casting experiments were carried out at utg using a Demag Technica continuous casting machine, type 30/10 D I MCP N. The experimental setup comprised three main components: a melting furnace, the casting tool, and a drawing unit. System controls for water flow,
temperature regulation, and kinematic control of the drawing process were developed by utg as previously described [23]. The temperature inside the mold is monitored by integrated thermocouples. The mold is cooled on both sides by copper coolers using water as the cooling medium. In the first stage of the casting tool, the Cu rod solidifies. Subsequently, Al melt is poured onto an additional ceramic and introduced atop the Cu in the second stage of the casting tool. Both Cu and Al rods are solidified and further cooled by the same cooling system. The resulting composite rod is extracted from the mold using rolling drawers that employ a go-stop withdrawal technique to avoid tearing the rod due to friction [24]. Each withdrawal cycle consisted of a 7.5 mm stroke followed by a waiting period of 1.3 s , resulting in a nominal casting speed of 263 mm/min in accordance with previous studies.

The upper side of the graphite mold at the Cu casting stage featured a rectangular structure designed to form the interface geometry within the Cu , resulting in grooves on its surface. However, the interface structure produced in the casting experiments was less pronounced than assumed in the rolling simulation. Due to this discrepancy, the comparability of the simulation results with the experiments is only possible to a limited extent.

Nevertheless, warm rolling of the compound rods was conducted at IBF using a universal rolling mill (Bühler VRW-400), configured in accordance with the setup assumed in the simulations. Three exemplary parameter sets of the simulations were selected for testing during the rolling trials:
a) Rolling at 300∘C with a speed ratio of 1.3 and a single pass achieving 50% height reduction.
b) Rolling at 300∘C with a speed ratio of 1.0 and achieving 50% height reduction in three passes according to schedule C (Table 1).
c) Rolling at 400∘C with a speed ratio of 1.0 and a single pass achieving 50% height reduction.

The rolled plates obtained from these experiments are shown in Figure 14. Consistent with the simulation results, none of the plates exhibited delamination at the Al/Cu layers. Setup a) produced a very homogeneous deformation between the Al and Cu layers, as reflected in both the visual inspection (Figure 14 a)) and the measurements of sample thickness and length (Table 2). In this setup, the thickness proportions of Al and Cu were nearly equal, and the elongation of Al was only 11% greater than that of Cu . In contrast, setup b) resulted in a plate where the Al layer exhibited 36% greater elongation compared to Cu and a 10% difference in thickness proportions. Both findings are in good agreement with simulation predictions, as in Figure 9 c) for setup a) and Figure 12 c) for setup b).

[picture]

Fig. 14. Experimental results of continuous compound casting and warm rolling with the setup a) - c).

Fig. 14. Experimental results of continuous compound casting and warm rolling with the setup a) - c).

Increasing the rolling temperature to 400∘C, as in setup c), did not lead to any directly observable changes in deformation behavior compared to setup a). However, based on the measured dimensions, the deformation appeared slightly less homogeneous than in setup a). Unlike the simulation results, the speed ratio appeared to influence the final curvature of the rolled plates. While setups b) and c) with speed ratio of 1.0 exhibited curvature values similar to those predicted by simulations (compare Figure 11), setup a) showed a significantly reduced curvature with a speed ratio of 1.3 .

Table 2. Dimensions of compound cast and warm rolled slabs for setup a) - c).

Parameter set | Proportion of Al layer thickness | Proportion of Cu layer thickness | Elongation of Al compared to Cu | Curvature of Cu in 1/mm
a) | 49% | 51% | 11% | 0.0039
b) | 45% | 55% | 36% | 0.0048
c) | 48% | 52% | 14% | 0.0047

Despite the discrepancies between the simulated and experimentally realized interface structures, several effects predicted by the simulations were already evident in the experimental outcomes. However, the experiments presented here should be regarded as preliminary findings. Further casting experiments employing controlled solidification fronts will be required in the future to achieve more distinct geometric structuring on the Cu surface and thereby improve mechanical interlocking with Al. Furthermore, to more conclusively evaluate the influence of parameters such as temperature on joint quality, a subsequent quantitative assessment of bond strength will be necessary.

Conclusion and Outlook

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

The obtained results of this study demonstrate that the introduction of a geometric interface can effectively compensate for shear stresses occurring during rolling and lead to a more homogeneous deformation of the layers. Compared to a flat interface, the mechanically interlocked layers endured higher shear stresses than expected with a flat interface. As the temperature is influencing the flow stress it was found to affect the strain distribution and the fluctuation of the shear stresses in the material compound significantly. The smallest difference in average strain between Al and Cu was observed at an initial temperature of 300∘C. Notably, this temperature also aligns with values reported in the literature as optimal for promoting strong metallurgical bonding while minimizing detrimental intermetallic compound formation [5,11]. However, comparably high fluctuations of shear stresses are expected which can be covered by the geometric interface. The influence on a possible metallurgical bond needs to be investigated. Adjusting the rolling speed ratio provided further control over deformation homogeneity. Increasing the speed ratio led to a reduction in average strain within the Al layer, bringing it closer to that of Cu and thus improving overall uniformity of deformation across the composite. However, no significant effect of rolling speed ratio on the compound material curvature was detected, suggesting that alternative strategies are required to achieve straightening during processing. When splitting the height reduction into several passes, the average strain of Al increased, resulting in less uniform elongation. All in all, a temperature of 300∘C, an increased speed ratio and large height reductions are suggested for good results.

Preliminary warm rolling experiments conducted at IBF using compound cast material produced at utg supported these findings. Although the interface structure formed on the copper surface were less pronounced than assumed in the simulations, several modeled effects were observed experimentally under selected conditions: Rolling at 300∘C with an increased speed ratio resulted in a homogeneous deformation between Al and Cu layers, while multiple passes led to larger elongation of the Al layer. Increasing the rolling temperature to 400∘C did not produce any directly observable changes in deformation behavior.

To fully evaluate joint quality, particularly bond strength, systematic mechanical testing will be required in future work. Further improvements to geometric structuring on copper surfaces will also be explored to enhance mechanical interlocking with aluminum. Finally, it is proposed to couple the current mechanical interlocking model with a model describing metallurgical bonding. To enable this approach, a bond strength model must be calibrated using experimental data obtained from laboratory-scale tensile and shear strength tests at different surface enlargements and temperatures as previously performed [25]. This combined modeling strategy will support more comprehensive optimization of process parameters.

Conflict of Interest

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

The authors declare that they have no conflict of interest.

Acknowledgment

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

This study was funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - Project-ID 457434681.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.55.pdf

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 69-76

doi: 10.4028/p-x7LK1s

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-17

Revised:
2026-02-13

Accepted:
2026-02-24

Online:
2026-04-15

A Hybrid Joining Strategy for Al-Cu Bimetallic Sheets Produced by Friction Stir Welding

The original version of this paper is available on https://www.scientific.net/MSF.1185.69.pdf

loannis Papantoniou 1,a* and Dimitrios Manolakos 1, b

1 School of Mechanical Engineering, Laboratory of Manufacturing Technology, National Technical

University of Athens, Athens, Greece

a*ipapanto@central.ntua.gr, b  manolako@central.ntua.gr

Abstract

The production of Al−Cu bimetallic sheets is of increasing interest for applications requiring a combination of lightweight performance and high thermal or electrical conductivity. Conventional fusion welding techniques are unsuitable due to excessive intermetallic compound (IMC) formation and poor bonding quality. Solid-state processes such as Friction Stir Welding (FSW) provide an attractive alternative; however, most studies aim to minimize heat input in order to suppress IMCs. In this work, a different approach is proposed. A hybrid joining strategy is employed, intentionally using controlled heat input and tool penetration to generate an extended stirring zone and a pronounced hook geometry. This results in mechanical interlocking combined with metallurgical bonding at the Al−Cu interface. Large-area bimetallic sheets were fabricated by FSW lap welding using four parallel passes to enlarge the bonded region. Microstructural characterization revealed the formation of continuous hook structures along the interface, promoting effective mechanical interlocking between the aluminum and copper layers. The integrity of the bimetallic sheets was further evaluated by cold rolling, which demonstrated excellent resistance to delamination despite local cracking of brittle IMCs. The results confirm that exploiting hook formation, rather than suppressing it, can provide a robust and scalable strategy for manufacturing Al−Cu bimetallic sheets by FSW.

Keywords: friction stir welding, Al−Cu, bimetallic sheets, hybrid joining strategy.

Introduction

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Aluminum-copper (Al-Cu) bimetallic sheets are attractive for lightweight structures and multifunctional components where low density must be combined with high electrical and thermal conductivity. However, producing a reliable Al−Cu interface remains challenging with fusion-based joining, because the thermal cycle promotes the formation of brittle Al−Cu intermetallic compounds (IMCs) and leads to weak or discontinuous bonding. As a result, solid-state processes, and particularly friction stir welding (FSW), have been widely investigated as an alternative route for dissimilar metal joining [1].

Despite these advantages, Al-Cu FSW remains governed by a fundamental trade-off. Increasing heat input and material mixing generally improves interfacial contact and bonding continuity, but at the same time accelerates IMC formation and increases the risk of local brittleness. Consequently, a significant portion of the existing literature focuses on optimizing process parameters in order to reduce heat input and suppress IMC formation, while maintaining acceptable mechanical performance [1−4]. In Al5083/Cu lap configurations, rotational and welding speeds have been shown to strongly influence material flow, hook morphology, IMC distribution, and lap-shear behavior [2,3]. Low heat input strategies for Al−Cu lap joints have also been proposed, demonstrating that sound joints can be achieved when IMC growth is limited [4].

At the same time, it has been increasingly recognized that the hook feature formed during lap FSW is not simply a geometric imperfection. Hook morphology directly affects the effective bonding area and the mechanical load transfer path, and therefore can either deteriorate or enhance joint integrity depending on its size, orientation, and continuity [5,6]. Nevertheless, in most multipass FSW studies, the hook is treated as a defect to be minimized or eliminated in subsequent passes in order to obtain a more uniform interface [5]. Other studies report that deeper tool penetration and increased stirring
can promote hook formation and increase bonded area, although this is often accompanied by increased IMC formation and associated cracking [6].

Motivated by these observations, the present work adopts a deliberately different manufacturing objective and joining philosophy. Rather than optimizing a single lap joint, the aim is to fabricate a large-area Al−Cu bimetallic sheet through the application of four parallel FSW passes, thereby enlarging the bonded interface in a controlled and scalable manner. Furthermore, instead of suppressing hook formation, a hybrid joining strategy is pursued. Controlled heat input and tool penetration are used to generate an extended stirring zone and a pronounced hook geometry, leading to mechanical interlocking combined with metallurgical bonding at the Al−Cu interface. The integrity of the resulting bimetallic sheet is finally evaluated through cold rolling, which acts as a stringent test of interfacial stability under severe plastic deformation.

Experimental Procedure

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Commercial AA5083 aluminum alloy plates and commercially pure copper plates were used for the fabrication of Al-Cu bimetallic sheets. Both materials had an initial thickness of 3 mm . The nominal chemical composition of the AA5083 alloy is presented in Table 1. The copper plates consisted of commercially pure copper with a minimum purity of 99.9wt.%, as specified by the supplier.

Table 1. Nominal chemical composition of AA5083 aluminum alloy (wt.%).

Element | Mg | Mn | Fe | Si | Cr | Cu | Zn | Ti | Al
wt.% | 4.0–4.9 | 0.4–1.0 | ≤0.40 | ≤0.40 | 0.05–0.25 | ≤0.10 | ≤0.25 | ≤0.15 | balance

Prior to welding, the faying surfaces were mechanically ground and degreased using ethanol in order to remove surface oxides and contaminants and ensure proper interfacial contact. Friction stir welding was performed using a conventional milling machine adapted for FSW operations. A schematic illustration of the process and tool geometry is presented in Fig. 1. The aluminum plate was positioned on the top side and the copper plate on the bottom side in a lap configuration. A nonconsumable tool made of tool steel was used. The tool consisted of a cylindrical shoulder with a diameter of 20 mm and a threaded cylindrical pin with a diameter of 5 mm , a pin height of 4 mm , and a thread pitch of 1 mm , as shown in Fig. 1. A tool tilt angle of 3∘ relative to the welding direction was applied in order to improve material consolidation and avoid defect formation.

[picture]

Fig. 1. Schematic illustration of the friction stir welding (FSW) process used to fabricate the Al−Cu bimetallic sheet, showing the rotating tool and the four parallel passes employed to produce a largearea bonded interface. The tool geometry and main dimensions are also indicated (shoulder diameter: 20 mm , pin diameter: 5 mm , pin height: 4 mm , thread pitch: 1 mm ).

Fig. 1. Schematic illustration of the friction stir welding (FSW) process used to fabricate the A l − C u bimetallic sheet, showing the rotating tool and the four parallel passes employed to produce a largearea bonded interface. The tool geometry and main dimensions are also indicated (shoulder diameter: 20 mm , pin diameter: 5 mm , pin height: 4 mm , thread pitch: 1 mm ).

The welding parameters were selected to provide the maximum achievable heat input for the specific FSW machine. The tool rotational speed was set to 1000 rpm , while the traverse speed was fixed at 13 mm/min, corresponding to the minimum discrete travel speed and the maximum rotational speed available. This parameter combination was intentionally chosen to promote extensive material softening, large stir zones, and the formation of a pronounced hook geometry at the Al−Cu interface.

To fabricate a large-area bimetallic sheet, four parallel FSW passes were applied on each specimen under identical process conditions. The distance between adjacent passes was fixed at 8 mm , as determined through a preliminary investigation. This spacing was intentionally selected to avoid overlap between adjacent stir zones, allowing the independent development of pronounced hook geometries in each pass while maintaining continuity of the Al−Cu bonded region across the sheet width.

After welding, transverse cross-sections were extracted perpendicular to the welding direction for microstructural characterization. Standard metallographic preparation procedures were followed, including grinding and polishing. Optical microscopy was used to examine the overall morphology of the stir zones, hook structures, and interface continuity. Scanning electron microscopy was employed to investigate local microstructural features and intermetallic compound formation at the Al−Cu interface.

Cold rolling experiments were conducted to assess the structural integrity of the bimetallic sheets under severe plastic deformation. Prior to rolling, the upper surface of the welded specimen was milled by removing approximately 0.5 mm of material in order to eliminate surface flash generated during FSW and to ensure a flat and clean surface condition. The appearance of the bimetallic specimen after surface cleaning and prior to rolling is shown in Fig. 2a. Rolling was performed using a laboratory rolling mill equipped with integrated force and torque sensors, allowing real-time monitoring of rolling conditions (Fig. 2b). After rolling, the specimens were examined for interfacial damage, crack initiation, and possible delamination between the aluminum and copper layers.

[picture]

Fig. 2. (a) Al-Cu bimetallic sheet after surface preparation by milling, prior to cold rolling, and (b) laboratory rolling mill used for deformation experiments, equipped with integrated force and torque sensors.

Fig. 2. (a) Al-Cu bimetallic sheet after surface preparation by milling, prior to cold rolling, and (b) laboratory rolling mill used for deformation experiments, equipped with integrated force and torque sensors.

Results and Discussion

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Macroscopic characterization and development of large-scale bimetallic sheets

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Macroscopic examination of the welded specimens confirms the successful fabrication of largescale Al−Cu bimetallic materials. As shown in Fig. 3, the application of four parallel friction stir welding passes resulted in a continuous bonded region extending across the sheet width, indicating that the adopted process strategy functions as a manufacturing route for bimetallic sheets rather than as a localized joining technique.

[picture]

Fig. 3. Schematic illustration (top) and corresponding cross-section (bottom) of the Al−Cu bimetallic sheet, showing mechanical interlocking through hook structures and metallurgical bonding at the interface.

Fig. 3. Schematic illustration (top) and corresponding cross-section (bottom) of the A l − C u bimetallic sheet, showing mechanical interlocking through hook structures and metallurgical bonding at the interface.

The macroscopic cross-sections reveal the formation of extensive stir zones, penetrating deeply into the copper layer. At the Al−Cu interface, pronounced and repetitive hook structures are clearly observed along the length of the welded region. The combined presence of large stir zones and welldeveloped hook geometries demonstrates that the selected high heat input conditions and tool penetration effectively promoted intense material flow and interfacial interaction.

The morphology of the hooks, together with their continuity across multiple passes, leads to a robust mechanical interlocking between the aluminum and copper layers. In parallel, the intimate contact generated by severe plastic deformation enables metallurgical bonding at the interface. Consequently, the Al−Cu bimetallic material is formed through a hybrid bonding mechanism, in which mechanical interlocking provides macroscopic structural integrity while metallurgical bonding contributes to interfacial cohesion.

Optical microscopy of the stir zone and hook morphology

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Optical microscopy observations provide further insight into the microstructural features governing the bonding mechanism. Representative micrographs taken from the left and right sides of the stir zone, as well as from the entire stirred region, are shown in Fig. 4a-c.

[picture]

Fig. 4. Optical micrographs of the stir zone showing hook structures and interfacial morphology: (a) left side of the stir zone, (b) right side of the stir zone, and (c) overall stir zone. Pronounced hook structures extending into the aluminum layer are observed, contributing to mechanical interlocking and metallurgical bonding. Localized interfacial defects are occasionally present at the base of the hook structures.

Fig. 4. Optical micrographs of the stir zone showing hook structures and interfacial morphology: (a) left side of the stir zone, (b) right side of the stir zone, and (c) overall stir zone. Pronounced hook structures extending into the aluminum layer are observed, contributing to mechanical interlocking and metallurgical bonding. Localized interfacial defects are occasionally present at the base of the hook structures.

In all cases, pronounced hook structures are observed to initiate at the Al−Cu interface and extend upward into the aluminum layer. The consistent upward growth of the hooks indicates intense plastic flow of copper into the aluminum during FSW, driven by the combined effects of tool rotation, traverse motion, and elevated temperature. The presence of well-developed hooks on both sides of the stir zone confirms the stability and repeatability of the material flow during welding.

Within the stir zone, the microstructure is dominated by an aluminum-rich matrix containing a significant population of copper-rich particles. These particles originate from the mechanical fragmentation and transport of copper into the aluminum during stirring. Their non-uniform but widespread distribution suggests efficient mechanical mixing, while also indicating the potential formation of Al−Cu intermetallic phases.

Localized defects, such as microcracks or small void-like features, are occasionally observed near the base of the hook structures. These regions are subjected to high strain localization and thermal exposure, making them susceptible to the formation and cracking of brittle phases. Importantly, these defects are isolated and discontinuous, and they do not form a continuous damage path along the interface. These localized interfacial defects are highlighted in Fig. 4 and remain spatially confined to limited regions near the hook roots, without extending along the interface or compromising the overall structural continuity of the bimetallic sheet.

Interfacial microstructure and metallurgical bonding

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Detailed examination of the Al−Cu interface was carried out using scanning electron microscopy combined with energy-dispersive X-ray spectroscopy. Figure 5 presents the SEM micrograph of the Al-Cu interface at high magnification, where a chemically and morphologically complex interfacial region is clearly observed.

[picture]

Fig. 5. SEM micrograph of the Al−Cu interface showing the formation of two distinct interfacial layers between the aluminum and copper substrates, attributed to aluminum-rich (CuAl2) and copperrich ( Cu9Al4 ) intermetallic compounds.

Fig. 5. SEM micrograph of the A l − C u interface showing the formation of two distinct interfacial layers between the aluminum and copper substrates, attributed to aluminum-rich ( C u A l 2 ) and copperrich ( C u 9 A l 4 ) intermetallic compounds.

As shown in Fig. 5, two distinct interfacial layers can be identified between the aluminum and copper substrates. The layer adjacent to the aluminum side exhibits a morphology consistent with aluminum-rich intermetallic compounds, while the layer adjacent to the copper side is enriched in copper. Based on their location, morphology, and contrast, these layers can be reasonably attributed to the θ-phase (CuAl2) near the aluminum side and the γ1-phase (Cu9Al4) near the copper side. The
formation of these phases is consistent with diffusion-controlled reactions promoted by severe plastic deformation and elevated temperature during friction stir welding.

To further investigate the chemical transition across the interface, an EDS line scan analysis was performed, as shown in Fig. 6. The line scan reveals a gradual decrease in aluminum concentration accompanied by a corresponding increase in copper concentration across the interfacial region. This compositional gradient confirms that the interface is not abrupt, but instead consists of a diffusiondriven transition zone.

The presence of two chemically distinct regions within the interfacial zone, highlighted in Fig. 6, provides additional evidence for the formation of layered Al−Cu intermetallic compounds. Such a configuration agrees well with previous reports on Al−Cu friction stir welded joints, where CuAl2 and Cu9Al4 layers are typically observed on the aluminum and copper sides, respectively [1,7]. Although these intermetallic compounds are inherently brittle, their thickness remain limited.

[picture]

Fig. 6. EDS line scan across the Al−Cu interface, illustrating the gradual transition from aluminumrich to copper-rich regions and the presence of two chemically distinct interfacial layers formed by diffusion during friction stir welding.

Fig. 6. EDS line scan across the A l − C u interface, illustrating the gradual transition from aluminumrich to copper-rich regions and the presence of two chemically distinct interfacial layers formed by diffusion during friction stir welding.

Interfacial integrity of the AI-Cu bimetallic sheet under cold rolling

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The mechanical integrity of the Al−Cu bimetallic sheet was further evaluated through successive cold rolling passes. Cold rolling imposes severe through-thickness plastic deformation and interfacial shear stresses, making it a stringent and widely accepted method for assessing the structural integrity and bonding quality of bimetallic materials. In insufficiently bonded systems, rolling typically promotes interfacial delamination or sliding between layers due to incompatible plastic flow and weak interfacial cohesion. Figure 7 presents the macroscopic appearance of the specimen after the 1st, 3rd, and 6th rolling passes, illustrating the progressive deformation imposed on the material.

[picture]

Fig. 7. Macroscopic evolution of the Al−Cu bimetallic sheet during cold rolling: (a) after the 1st rolling pass, (b) after the 3rd rolling pass, and (c) after the 6th rolling pass, demonstrating progressive deformation without macroscopic delamination of the Al−Cu interface.

Fig. 7. Macroscopic evolution of the A l − C u bimetallic sheet during cold rolling: (a) after the 1st rolling pass, (b) after the 3rd rolling pass, and (c) after the 6th rolling pass, demonstrating progressive deformation without macroscopic delamination of the A l − C u interface.

As shown in Fig. 7, the bimetallic sheet undergoes increasing levels of deformation with each rolling pass. Despite the severe plastic deformation introduced during rolling, no macroscopic delamination or separation between the aluminum and copper layers is observed, even after the 6th rolling pass. This behavior indicates that the Al−Cu interface remains mechanically stable under substantial through-thickness strain.

The evolution of the average rolling force as a function of pass number, shown in Fig. 8, further supports this observation. The rolling force increases progressively with increasing deformation, reaching values above 150 kN in the final pass. The rolling force was continuously recorded during each rolling pass, and the values presented in Fig. 8 correspond to the average force calculated from these continuous measurements. Importantly, no abrupt drops or instabilities were observed in the force evolution during rolling, which would otherwise indicate interfacial failure or sliding between the bonded layers. The smooth and continuous evolution of rolling force suggests effective load transfer across the Al−Cu interface and confirms that the bimetallic sheet behaves as a single structural entity during deformation.

The rolling force generally increases with successive passes due to strain hardening and the associated increase in deformation resistance, as well as the slight increase in specimen width during rolling, which increases the effective contact area. The force values recorded for passes 3 to 5 are very close, indicating a relatively stable mechanical response. The slight decrease observed in passes 4 and 5 may be attributed to localized fracture and redistribution of brittle intermetallic phases within the stir zone, which can locally reduce deformation resistance. However, the overall trend of increasing rolling force confirms progressive work hardening and effective load transfer across the Al−Cu interface.

[picture]

Fig. 8. Average rolling force versus rolling pass number for the Al−Cu bimetallic sheet, indicating stable load transfer during successive deformation passes without interfacial failure.

Fig. 8. Average rolling force versus rolling pass number for the A l − C u bimetallic sheet, indicating stable load transfer during successive deformation passes without interfacial failure.

The absence of macroscopic delamination during rolling can be directly attributed to the hybrid bonding mechanism developed during friction stir welding. The pronounced hook structures generated at the Al−Cu interface provide strong mechanical interlocking, which resists interfacial sliding and separation under compressive stresses. At the same time, metallurgical bonding and localized intermetallic compound formation enhance interfacial cohesion, contributing to the overall stability of the joint.

Although localized cracking of brittle intermetallic phases was observed at the microscale, these features do not propagate into continuous interfacial damage under rolling. Instead, the mechanical interlocking mechanism dominates the macroscopic deformation response, ensuring the preservation of interfacial integrity throughout successive rolling passes.

Overall, the rolling experiments demonstrate that the fabricated Al−Cu bimetallic sheet exhibits high resistance to interfacial failure under severe plastic deformation. This behavior confirms the effectiveness of the proposed joining strategy for the production of mechanically robust bimetallic sheets suitable for subsequent forming operations.

Conclusion

The original version of this paper is available on https://www.scientific.net/MSF.1185.69.pdf

In this work, a hybrid joining strategy was employed for the fabrication of Al-Cu bimetallic sheets using friction stir welding. Instead of suppressing hook formation, controlled heat input and tool penetration were deliberately selected to promote the development of large stir zones and pronounced hook geometries, leading to mechanical interlocking combined with metallurgical bonding at the AlCu interface.

Macroscopic and microscopic analyses confirmed the successful formation of large-scale bimetallic materials through multiple parallel FSW passes. Optical and SEM observations revealed extensive hook structures extending into the aluminum layer, accompanied by localized diffusiondriven intermetallic compound formation at the interface. These microstructural features were shown to be mechanically effective, as confirmed by cold rolling experiments used as a stringent validation of interfacial integrity. The bimetallic sheets withstood severe plastic deformation without macroscopic delamination, even after multiple rolling passes. The progressive increase in rolling force and the absence of load instabilities further confirmed effective load transfer across the Al−Cu interface and the structural coherence of the bimetallic sheet.

Overall, the results demonstrate that exploiting mechanical interlocking in combination with metallurgical bonding provides an effective and scalable route for producing mechanically robust AlCu bimetallic sheets. The proposed strategy demonstrates a high tolerance to severe plastic deformation, maintaining interfacial integrity without complete delamination despite localized fracture of brittle intermetallic phases. Such damage-tolerant Al−Cu bimetallic systems are particularly relevant for structural and functional applications where resistance to catastrophic delamination under severe deformation or impact-like loading is required, rather than preservation of intermetallic integrity.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.69.pdf

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P. Xue, B.L. Xiao, D. Wang, Z.Y. Ma, Achieving high property friction stir welded aluminium/copper lap joint at low heat input, Sci. Technol. Weld. Join. 16 (2011) 657-661. .

Y. Wei, F. Sun, H. Li, X. Peng, J. Zou, The microstructure and mechanical properties of multipass friction stir-welded lap joints and interface transfer during butt joining of thick plates, Adv. Eng. Mater. 22 (2020) e202000346. . 202000346.

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 77-87

doi: 10.4028/p-hOTSU9

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-18

Revised:
2026-02-10

Accepted:
2026-02-24

Online:
2026-04-15

Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

PIWEK, Armin 1,a*, PEDDINGHAUS, Julius 1, b,UHE Johanna 1,c, BRUNOTTE Kai 1, d and BEHRENS Bernd-Arno 1,e

1 Institute of Forming Technology and Machines, Leibniz University Hannover, An der Universität 2, 30823 Garbsen

a* piwek@ifum.uni-hannover.de, b  peddinghaus@ifum.uni-hannover.de, c uhe@ifum.uni-hannover.de, d brunotte@ifum.uni-hannover.de, e behrens@ifum.uni-hannover.de (*corresponding author)

Keywords: Joining, Hybrid Materials, Heat Superimposition, Lightweight.

Abstract

Hybrid joints made of steel and aluminium alloy produced by rotary friction welding enable load-adapted lightweight components. However, a major challenge is the inhomogeneous radial temperature distribution caused by different relative velocities between the specimen centre and edge during rotation. This effect leads to local insufficient bonding and reduces the overall joint strength, especially in the centre, where low relative rotation speeds occur. Previous studies mainly addressed preheating before the friction phase, whereas superimposed heating during the upsetting phase has not been investigated so far. To achieve temperature equalisation along the cross-section during rotary friction welding, a modified KUKA Genius plus machine equipped with joule heating was used to introduce an electric current during the upsetting phase. Experiments were conducted on EN AW6082 (AA-6082) joined to 20 MnCr 5 (AISI 5120 H ). A three-step variation of current intensity (10, 24 and 36 A/mm2 ), alongside a reference without current, was investigated. Temperatures were monitored using type K thermocouples, confirming temperature equalisation. Mechanical performance was assessed by uniaxial tensile tests, while hardness measurements and metallographic analyses characterised the influence of superimposed heating on the interfacial microstructure. Joint strength improves up to 17% with increasing current, even under otherwise unsuitable welding parameters that would normally result in insufficient bond strength. This improvement is linked to a uniform temperature distribution and enhanced material flow, resulting in a defect-free specimen centre.

1. Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

Increasing demands for energy-efficient and resource-conserving structures drive the development of hybrid lightweight components that combine high specific strength with reduced mass. Material systems such as steel-aluminium pairs offer substantial potential for functional integration and weight reduction, making them attractive for industrial applications ranging from automotive structures to production technology [ 1,2 ].

Reliable and microstructurally controlled joining processes are essential, yet conventional fusionbased techniques frequently suffer from brittle intermetallic compound formation and inadequate suitability for dissimilar material pairings. Solid-state processes, particularly rotary friction welding (RFW), therefore offer a promising alternative due to lower heat input. Nevertheless, a major limitation arises from the inherent radial temperature gradient due to the increasing tangential velocity along the radius. With a maximum at the outer radius and a minimum of zero at the centre, uneven frictional heat generation develops on the welding surfaces. Consequently, different plastic strains and thus grain sizes occur along the radius in the subsequent upsetting phase, which can impair the homogeneity and strength of the joint [ 3,4,5 ]. To ensure sufficient temperatures and material flow, an increase in friction time and friction pressure is required. Efforts to mitigate radial temperature gradient have so far focused mainly before, after or in the friction phase, e. g., by adjusting process parameters [2,6] or preheating the components [7,8,9]. The required thermal energy is provided by
an external heat input rather than by frictional heating alone. This approach drastically reduces the necessary upsetting force and friction time, resulting in minimal flash formation, reduced machine dimensions and smaller heat-affected zones compared to conventional friction welding [10, 11].

In addition to the ongoing development of methods aimed at achieving more homogeneous heat distribution, it appears promising to shift the research focus more strongly towards in-process concepts rather than purely pre- or post-weld treatments. Such strategies offer potential to meet both technological and economic requirements in industrial applications. Therefore, the present study provides the first systematic investigation of current assisted heating during the upsetting phase in RFW for the material pairing 20 MnCr 5 and EN AW-6082. The objective is to determine how an additional joule heating affects radial temperature equalisation, interfacial microstructure, specimen shortening and the tensile strength of the bond. The findings aim to deepen the understanding of thermomechanical interactions within the process and to identify industrially relevant potentials for more robust and functionally optimised hybrid lightweight joints in regards of material utilisation, mechanical load requirements, and overall process controllability.

2. State of the Art

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

2.1 Temperature Distribution in Rotary Friction Welding

RFW is a solid-state joining process in which two components are brought into axial contact, while one workpiece undergoes relative motion so that frictional heat is generated and the material at the interface becomes locally plasticised. Without the formation of a melt layer, the joining region is deformed and recrystallised, while the subsequent upsetting phase expels oxides as well as contaminants into the characteristic flash and establishes a sound metallurgical bond [5, 12]. Numerous studies confirm the suitability of RFW for both similar and dissimilar metallic combinations, as the short thermal cycles limit the growth of brittle intermetallic compounds and enable the formation of refined interface microstructures [1, 13]. Detailed microstructural analyses further indicate dynamically recrystallised grains and strong gradients in deformation and phase evolution within the thermo-mechanically affected zone[5]. However, the process also entails inherent challenges. Insufficient or uneven temperature development can hinder optimal plasticisation and promote the formation of critical intermetallic layers in dissimilar joints, particularly in aluminium-steel systems [1,13]. These thermal limitations are a central motivation for recent studies exploring the use of additional external heating to homogenise the thermal boundary conditions and improve process stability and dissimilar joint performance.

A homogeneous and well-controlled temperature field is essential in RFW, since spatial thermal gradients govern local softening, interfacial shear conditions and the formation and growth of reaction layers. Several studies have demonstrated an inherent radially non-uniform thermo-mechanical field in RFW, as frictional heat generation increases with radius due to the linear rise in relative velocity. Experimental work by Alves et al. on aluminium-steel confirmed the strong sensitivity of the joint to such temperature gradients, revealing incomplete plasticisation and reduced strength under insufficient heat input, while excessive thermal exposure promotes pronounced softening and material displacement [14]. Extending these findings, Mishra et al. identified radial temperature gradients as a key factor driving radially heterogeneous plastic flow and interfacial deformation, particularly promoting weak bonding or defect formation in the joint centre [15]. In a similar study, MA ET AL. observed a reduction in strength for RFW joints between stainless steel X5CrNi18-10 and aluminium EN AW-2014 under both insufficient and excessive heat input. Insufficient thermal activation limited interfacial bonding, whereas excessive heat input promoted uneven and excessive growth of intermetallic compounds (IMC). Conclusively, only a moderate and radially balanced temperature field enabled the formation of a thin, continuous reaction layer and thus ensures a globally reliable joint [5]. Based on the presented studies, the aim is therefore to align the inhomogeneous heat distribution occurring in RFW with the more uniform temperature distributions prevailing in linear or orbital friction welding, since these processes yield a significantly more homogeneous heat input [4, 16]. This could not only promote the formation of more uniform
microstructural zones, but also contribute to improved and more consistent stress states in the upsetting phase, which would again increase joint strength.

2.2 Additional Heating in Rotary Friction Welding

Various approaches are currently being investigated in which the process is supplemented by an additional heat source. Many current studies focus on preheating one or both joining partners to ensure a more uniform temperature distribution in the joining zone. Within this context, Dahlan et AL. examined the effect of targeted preheating in combination with varying friction times on the joint strength of RFW between stainless steel X5CrNi18-10 and pure zinc semi-finished products. The steel component was heated to different target temperatures to reduce the effective friction time. Based on tensile tests, an increase in strength with a moderate preheating temperature of 100∘C could be observed. However, increasing the preheating temperature beyond this moderate level led to a reduction in tensile strength, which was attributed to increased enrichment of Fe,Cr,O, and C elements in the joining interface [17]. A similar conclusion was drawn by Haibin et al., who analysed preheating of the H59 brass component during RFW with EN AW-1100 aluminium to improve plastic deformation on the brass side while simultaneously controlling the formation of IMC. For this purpose, the brass part was heated to two different preheating temperatures by inductive heating immediately before welding. The highest joint strength was obtained at a temperature of 327∘C, which was attributed to the formation of a uniform, thin IMC of 3μ m. In contrast, preheating to 600∘C resulted in an excessively thick IMC of 16μ m and promoted thermally induced cracking, significantly reducing joint strength [18]. Beyond inductive preheating, alternative heat input strategies have also been explored. Mullo ET Al. investigated RFW between C45E steel and EN AW-2017-T4 aluminium while applying laser-based preheating to the steel component. In this case, increasing laser power consistently correlated with higher tensile strength, demonstrating that locally applied and controllable heat sources can effectively enhance joint performance [7].

Closely related to these concepts, the so-called low-force friction welding process has been developed by Manufacturing Technology, Inc. in collaboration with Edison Welding Institute. In this case, targeted preheating of the joining partners primarily serves to significantly reduce the process forces required during welding. In addition to induction heating, resistance-based external heating sources have also been tested for this process [19]. Apart from reduced process forces, this approach offers further advantages, including smaller flash formation, the possibility of using more compact machines, increased material compatibility, and shorter cycle times due to reduced friction time [10, 11].

Analogous to the upsetting phase of RFW, LIM ET AL. investigated a solid-state joining process based on a combination of Joule heating and mechanical pressure. The applied pressure was selected to match the intersection of the temperature-dependent yield strengths of the two materials Ti-6Al4 V titanium alloy and SUS316L stainless steel. In this way, uniform plastic deformation of both joining partners along the entire joint line could be achieved. The experimental results demonstrate that this process produces a significantly more homogeneous temperature distribution in the joining area compared with RFW. Notably, the temperature at the centre of the joining zone was slightly higher, which was attributed to increased cooling by ambient air in the peripheral region [20].

The reviewed studies clearly demonstrate the potentials of supplementing solid state welding with additional heat input to adjust the thermo-mechanical conditions in the joining zone. Preheating has been shown to improve joint strength and reduce process forces if applied within a narrow thermal window. However, excessive or poorly controlled heat input may lead to detrimental effects such as excessive IMC growth, microstructural degradation, or thermally induced cracking. Moreover, conventional preheating approaches often act globally on one joining partner and remain decoupled from the actual heat generation during the friction and upsetting phases. Therefore, the present work investigates a current-assisted RFW approach, in which additional heat is introduced directly into the joining zone by joule heating during upsetting. The objective is to perform a combined microstructural and mechanical evaluation of steel-aluminium joints, focusing on how different levels of electrical current modify the local thermo-mechanical conditions, interfacial microstructure, deformation behaviour, and resulting joint strength.

3. Materials and Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

For the investigation, the case-hardening steel 20 MnCr (AISI 5120H/1.7147 ) in the soft-annealed condition (hardness 180 HV1) as well as the heat-treatable aluminium alloy EN AW-6082 (AA-6082) in the T6 condition (hardness 100 HV1) were used. The chemical composition was measured by spark emission spectroscopy and is shown in Tab. 1.

Table 1. Chemical composition (mean value of 5)

| Si | Mn | Mg | C | Cr | Fe | Al
EN AW-6082 | 1.08 | 0.45 | 0.83 | - | 0.02 | 0.40 | 96.95
20MnCr5 | 0.24 | 1.29 | - | 0.17 | 1.07 | 97.04 | 0.02

The two materials were joined as cylindrical rods with a diameter D=40 mm and a length l=100 mm by means of RFW on a KUKA Genius Plus system (see Fig. 1 a). The friction welding setup includes a conductive heating unit consisting of copper conductors integrated into the clamping systems. During the upsetting phase, the specimens can be heated by a constant electric DC heating current based on the principle of resistance heating once the components are in contact. When the current is applied, the joining zone is additionally heated by the contact resistance at the faying surfaces of the semi-finished parts, in addition to the frictional heat introduced beforehand (see Fig. 1 b).

[picture]

Fig. 1. a) Experimental setup of RFW with and without additional heating, b) exemplary RFW process with different current densities ( n=700 min−1,pF=100MPa,tF=0.2 s,pU=225MPa ).

Fig. 1. a) Experimental setup of RFW with and without additional heating, b) exemplary RFW process with different current densities ( n = 700 m i n − 1 , p F = 100 M P a , t F = 0.2 s , p U = 225 M P a ).

At the beginning of the process, the clamped steel part is accelerated to a rotational speed of n=700 min−1. In the beginning of the friction phase, the specimens are brought into contact for 1 s with an axial force of 35 kN . This is followed by the initial friction phase at a pressure of 100 MPa , during which the heat required for the process is generated. After the friction time ( tF ), the upsetting phase follows, in which an increased pressure ( p∪ ) is applied. Different current densities are set at the beginning of the upsetting phase and are varied between 0 A/mm2,10 A/mm2×4 s,24 A/mm2×2 s and 36 A/mm2×1.5 s for different friction times ( 0.1 s;0.2 s ) and upsetting pressures ( 150 MPa or 225 MPa ). The electric current was applied exclusively during the upsetting phase, as this is the stage at which full interfacial contact is established under high axial pressure and relative motion has ceased. Under these conditions, resistance heating can be introduced in a controlled and localised manner directly at the joining interface. Applying current earlier would interfere with frictional heat generation and could promote premature softening of the aluminium alloy, leading to unstable process conditions and increased flash formation.

The experiments were carried out using a full factorial experimental design with twelve parameter sets using specimen shortening and tensile strength as a response variable. Uniaxial tensile testing was conducted to determine the bond strength using a universal testing machine of the AllroundLine Z250 type (ZwickRoell). The specimen geometry corresponds to a round tensile test specimen designed in accordance with DIN EN ISO 6892-1, with a gauge diameter of 30 mm and a gauge length of 85 mm . The reduction of the diameter is necessary to ensure fracture within this testing
method at a maximum test force of 250 kN . The comparability between the specimen series is not limited due to the identical machining. Furthermore, to evaluate the radial temperature distribution, temperatures were measured at a representative current density of 24 A/mm2 using type K thermocouples positioned in the centre axis and at the periphery, approximately 6 mm from the joining zone, enabling a direct comparison between the conventional and the current-assisted process. Furthermore, to assess the influence of the electric current on the resulting microstructure, the specimens were characterised by metallographic analysis using light optical microscopy and hardness measurements.

4. Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

4.1 Temperature Distribution

The temperature profiles presented in Fig. 2 a indicate a significant modification of the thermal behaviour when applying electric current in the joining process. For identical RFW parameters, the application of a current density of 24 A/mm2 increases the absolute temperature level, with peak temperatures of 200∘C reached compared to 150−160∘C in the conventional process. The apparent time delay in the temperature development is caused by transient heat conduction from the joining interface to the measurement positions. Since the thermocouples are located 6 mm from the weld interface, the recorded temperature reflects the thermal response of the surrounding material rather than the instantaneous interface temperature. As a result, temperature maxima occur with a temporal offset after changes in heat generation, particularly at the transition from the friction to the upsetting phase.

Without current assistance, the temperature reaches its maximum at the beginning of the upsetting phase and subsequently decreases. The lower relative velocities in the central region lead to delayed heat generation and, consequently, to pronounced radial temperature gradients, which are expected to be even more severe directly at the weld interface. With the introduction of electric current, additional heating in the joining zone is generated, effectively reducing the temperature gradient between the centre and the edge area, see Fig. 2 b. While the conventional process exhibits a maximum temperature difference of 56 K between centre and periphery, this gradient is reduced to about 16 K when electric current is applied, indicating a substantially more homogeneous thermal field.

[picture]

Fig. 2. Exemplary temperature curve, 6 mm distant to joining zone: a) measurements of the centre and edge area, b) temperature difference ΔT between centre and edge ( n=700 min−1,pF=100MPa, tF=0.2 s,pU=225MPa ).

Fig. 2. Exemplary temperature curve, 6 mm distant to joining zone: a) measurements of the centre and edge area, b) temperature difference Δ T between centre and edge ( n = 700 m i n − 1 , p F = 100 M P a , t F = 0.2 s , p U = 225 M P a ).

4.2 Flash Formation

Macroscopic cross-sections of joints between 20 MnCr 5 steel and EN AW-6082 aluminium produced at different current densities are shown in Fig. 3. In all cases, flash formation is predominantly observed on the aluminium side, while the steel component remains largely undeformed. The aluminium flash exhibits a characteristic curved and folded geometry, indicating pronounced plastic flow during the friction and upsetting phases. As a result of the upsetting pressure and the differing flow stresses depending on temperature, flash formation occurs, and the material displaced from the joining zone leads to a shortening of the semi-finished part, which should be avoided for both ecological and economic reasons.

With increasing current density, the flash becomes more pronounced. The observed flash formation is governed by the strongly temperature-dependent flow behaviour of the aluminium alloy. Compared to steel, EN AW-6082 exhibits significantly lower flow stresses, which decrease rapidly with increasing temperature [21]. While the steel component remains mainly unformed due to its higher warm strength, the aluminium becomes highly plasticised in the joining zone and is displaced outward under the applied axial load. The application of electric current introduces additional resistance heating, increasing the local temperature beyond that generated by friction alone, as seen in Fig. 3.

Identical welding parameters were applied for all current densities. Therefore, the increased flash formation is primarily a consequence of the additional heat input provided by the electric current. Therefore, reduced friction times and lower axial forces may be applied while maintaining a sufficient plasticisation and flash formation. Overall, the superimposed heating expands the usable process window, allowing greater flexibility in parameter selection.

[picture]

Fig. 3. Flash formation depending on current density, same RFW parameters ( n=700 min−1, pF=100MPa,tF=0,2 s,pU=225MPa,tU=6 s ).

Fig. 3. Flash formation depending on current density, same RFW parameters ( n = 700 m i n − 1 , p F = 100 M P a , t F = 0 , 2 s , p U = 225 M P a , t U = 6 s ).

4.3 Bond Strength and Shortening

To evaluate the effect of electric current on the bond strength, the RFW specimens were investigated by uniaxial tensile testing, and the resulting bond strength was correlated with the process-induced shortening, as shown in Fig. 4 a and Fig. 4 b. Overall, a beneficial influence of current assistance on the bond strength is evident.

[picture]

Fig. 4. a) Bond strength, b) specimen shortening, ( 3 measurements per parameter set).

Fig. 4. a) Bond strength, b) specimen shortening, ( 3 measurements per parameter set).

Moderate to high current densities lead to a pronounced increase in bond strength. In particular, a density of 24 A/mm2 provides a favorable balance between high average strength (up to 273 MPa ) and similar material loss compared to 0 A/mm2, making it advantageous from both a resourceefficiency and economic perspective. Although a higher current density of 36 A/mm2 enable high strengths of up to 273 MPa for short friction times, this advantage is diminished by the increased shortening if high upsetting pressures are present. However, at lower upsetting pressures of 150 MPa , current assistance ( 36 A/mm2 ) allows a noticeable improvement up to 17% from 228 MPa to 266 MPa while maintaining a similar level of shortening compared to the reference. Therefore, the combination of reduced friction time and additional electrical heating effectively expands the usable process window, allowing mechanically less severe parameter combinations to achieve sufficient bonds. In contrast, at extended friction times of 0.2 s , a detrimental effect is observed for low current density ( 10 A/mm2 ). This highlights that current assistance does not universally improve bond quality, but must be carefully matched with the friction time and upsetting pressure to achieve a beneficial thermal balance at a reduced material loss.

4.4 Metallography

4.4.1 Microstructural Characterisation

Pronounced differences in the microstructure are observed as a function of radial position and applied current density, as shown in Fig. 5.

[picture]

Fig. 5. Micrographs dependent on position and current density. Steel microstructure: a) - d) edge and e) - h) centre, i) -l) central microstructure in the aluminium ( n=700 min−1,pF=100MPa,tF=0,2 s, pU=225MPa,tU=6 s ).

Fig. 5. Micrographs dependent on position and current density. Steel microstructure: a) - d) edge and e) - h) centre, i) -l) central microstructure in the aluminium ( n = 700 m i n − 1 , p F = 100 M P a , t F = 0 , 2 s , p U = 225 M P a , t U = 6 s ).

In the edge region of the joint (Fig. 5 a-d), where frictional heat generation is highest, a continuous bonding interface is formed for all conditions. The steel exhibits a deformed ferritic-pearlitic microstructure adjacent to the joining line, with increasing waviness and plasticisation at higher current densities. In contrast, the centre of the joining zone (Fig. 5 e-h) indicate a strong dependence on current assistance. In the reference condition with 0 A/mm2 (Fig. 5𝐞 ), the central interface is locally discontinuous, and regions of insufficient material flow are observed. The microstructural observations are characteristic of insufficient thermomechanical activation in the centre, where frictional heating is reduced, and align with central bonding defects reported in previous studies [1, 3, 15]. The steel microstructure appears strongly deformed but inhomogeneous, reflecting nonuniform local strain and temperature conditions rather than systematic grain coarsening or phase transformation effects. At a current density of 10 A/mm2 (Fig. 5 b, f), the interfacial continuity is slightly improved compared to the reference condition. However, the centre still exhibits noticeable microstructural gradients, and localised flow deficiencies remain visible. The additional heat input at this current density is therefore insufficient to fully compensate for the reduced frictional heating in the centre.

At a current density of 24 A/mm2 (Fig. 5 c, g), a continuous interface is formed across the entire joint width. The steel microstructure adjacent to the interface appears more uniformly plastified from edge to centre, and radial variations in interfacial morphology are reduced. This homogenisation indicates a more uniform thermal and deformation state, resulting from the combined action of mechanical friction and joule heating. The improved interfacial continuity is in good agreement with the enhanced bond strength observed for this condition.

At the highest current density of 36 A/mm2 (Fig. 5𝐝,𝐡 ), the interface remains continuous and uniform throughout the joint. Plastic deformation in the steel near the joining zone is present due to elevated temperatures and increased material flow. While no interfacial defects are observed, the microstructural appearance suggests an increased dominance of thermal effects, which correlates with the experimentally observed increase in axial shortening at higher upsetting forces.

The corresponding aluminium microstructures adjacent to the joining zone (Fig. 5 i-l) further support these findings. In the reference condition (Fig.5i), the aluminium exhibits oriented deformation bands, indicative of shear deformation under limited thermal softening. With increasing current density (Fig. 5𝐣,𝐤 ), the deformation structures become progressively more diffuse, and orientation gradients are reduced. At 24 A/mm2 and 36 A/mm2, the aluminium microstructure appears more homogeneous, reflecting enhanced thermal activation and a more uniform strain distribution near the interface.

Overall, based on the metallographic analysis, current-assisted RFW effectively reduces radial microstructural and interfacial gradients, particularly in the joint centre. Moderate current densities promote a homogeneous thermomechanical state in both joining partners, improving bonding conditions and hence suppressing central bonding defects. At excessively high current densities, microstructural homogenisation is maintained, but at the expense of increased thermal exposure, highlighting the importance of an optimised balance between electrical and mechanical heat input.

4.4.2 Hardness Measurements

The steel-side hardness profiles remain on a comparable level for the reference condition ( 0 A/mm2 ) and current densities up to 24 A/mm2, with typical values of 175-200 HV in near the joining zone, see Fig. 6. At a current density of 36 A/mm2, a distinct hardness increase is observed, with local values reaching 200−220HV. This increase cannot be explained by grain refinement, as metallographic analysis reveals local grain coarsening near the joining zone, see Fig. 5. Consequently, the classical Hall-Petch relationship is not the dominant strengthening mechanism in this case. Instead, the hardness increase is governed by deformation-induced strengthening, including elevated dislocation density and a strongly deformed ferrite-pearlite morphology resulting from intensified plastic flow during welding and upsetting. Such deformation-related effects are known to outweigh grain-size-related softening when recovery and recrystallisation remain incomplete [22].

On the aluminium side, a pronounced hardness minimum is observed near the joining zone for all conditions, with values decreasing to 75 HV , followed by a gradual recovery towards the base material. At higher current densities, this local softening becomes slightly less pronounced, reflecting increased thermal exposure and plastic deformation, which is consistent with the observed material flow and flash formation. During upsetting, the additional heat input leads to a pronounced reduction in the flow stress of aluminium, promoting intensified radial material flow both in the joint centre and edge area. Thermally softened aluminium is preferentially extruded outward during upsetting, forming an enlarged flash and removing highly heated material from the interface [ 1,3,15 ].

[picture]

Fig. 6. Hardness measurements. a) Hardness in the centre, b) hardness 10 mm distant from the edge. ( n=700 min−1,pF=100MPa,tF=0,2 s,pU=225MPa,tU=6 s ).

Fig. 6. Hardness measurements. a) Hardness in the centre, b) hardness 10 mm distant from the edge. ( n = 700 m i n − 1 , p F = 100 M P a , t F = 0 , 2 s , p U = 225 M P a , t U = 6 s ).

5. Summary

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

The application of electric current leads to a homogenisation of the temperature distribution between the centre and the peripheral regions. With increasing current density, the bond strength further improves up to 17% compared to the reference condition. Moreover, the use of electric current reduces the axial force required during the upsetting phase to achieve a high-strength joint. This reduction in the upsetting force required potentially allows larger cross-sections to be joined without the need to increase the mechanical capacity of the machine. Although shortening increases with current application for identical process parameters, an appropriate adjustment of the process parameters allows a reduction in material loss and cycle time. As a result, parameter sets that would be unsuitable without current assistance can be potentially applied.

6. Outlook

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

For future work, local tensile tests at different radial positions are planned to assess the local joint quality along the radius in greater detail. In addition, the influence of current application on the radial distribution of intermetallic compounds will be investigated. A further outlook includes the welding of materials that are conventionally considered unsuitable due to mismatched thermal expansion behaviour, which otherwise prevents the formation of a sufficient joint. Finally, with regard to economic and environmental aspects, a comparative analysis of electrical energy consumption, process efficiency, and potential material savings is envisaged. Based on these factors, the associated CO2-emission reduction potential of the respective processes can be quantified.

Acknowledgement

The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

Funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - CRC 1153 subproject B03-252662854. The authors thank KUKA Deutschland GmbH for providing the conductive heating unit.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.77.pdf

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U. Dilthey, Schweißtechnische Fertigungsverfahren 1: Schweiß- und Schneidtechnologien, Springer-Verlag, Berlin / Heidelberg, 2006.

S. Herbst, H. J. Maier, F. Nürnberger, Strategies for the Heat Treatment of Steel-Aluminium Hybrid Components, HTM J. Heat Treatm. Mat. 73 (2018) 5, 268-282,

E.P. Alves, F.P. Neto, C.Y. An, E.C. da Silva, Experimental determination of temperature during rotary friction welding of AA1050 aluminum with AISI 304 stainless steel, J. Aerosp. Technol. Manag. 4 (2012) 61-67.

N. K. Mishra, A. K. Singh, S. K. Panda et al., Interfacial inhomogeneous plastic deformation during rotary friction welding of dissimilar materials, J. Adv. Join. Process. 10 (2024) 100245.

I. Martinek, H.Goldau, und F. Trommer, "Grundsatzuntersuchungen zum Orbitalreibschweißen des unlegierten Baustahls S355J2," Innovationspreis Reibschweißen, 2010.[Online]. Available: _trommer.pdf

H. Dahlan, A.K. Nasution, S.A. Zuhdi, M. Rusli, Study of the effect of friction time and preheating on the joint mechanical properties of friction welded SS316-pure Zn, Appl. Sci. 13 (2023) 988.

G. Haibin, X. Guifeng, L. Jingcheng, L. Hao, Rotary friction welding of pure aluminum to preheated brass, Weld. World 66 (2022) 2371-2376.

Z. Roelofs, J. E. Gould, Preliminary results on low force friction welds of dissimilar materials: Aluminum 6061 to 1020 steel tubes. In: Low Force Friction Welding of Bimetallics, Manufacturing Technology, Inc., 2020.

Y. Lim, Y. Morisada, H. Liu, und H. Fujii, "Ti-6Al-4V/SUS316L dissimilar joints with ultrahigh joint efficiency fabricated by a novel pressure-controlled joule heat forge welding method," Journal of Materials Processing Technology, vol. 298, p. 117283, 2021,

F. J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, 2nd ed., Elsevier, Oxford, 2004.

I. Schindler, P. Kawulok, V. Očenášek, et al. Flow Stress and Hot Deformation Activation Energy of 6082 Aluminium Alloy Influenced by Initial Structural State Metals 9 (2019) 1248.

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 89-97

doi: 10.4028/p-Fx9Ipm

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-18

Revised:
2026-02-16

Accepted:
2026-02-24

Online:
2026-04-15

Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

Riccardo Puleo 1,a*, Salvatore Russo 1,b, Gianluca Buffa 1,c and Livan Fratini 1,d

1 Department of Engineering, University of Palermo, Viale delle Scienze, Palermo 90128, Italy

a* riccardo.puleo01@unipa.it, b salvatore.russo18@unipa.it, c gianluca.buffa@unipa.it, divan.fratini@unipa.it

Keywords: FSE, aluminium, recycling, extrusion, solid state.

Abstract

Nowadays, the growing demand for sustainable solutions in manufacturing has shifted research attention toward innovative recycling strategies. Among these, the Solid-State Recycling (SSR) technique has emerged as a viable approach to transform metal swarf into new products. Within the SSR family, friction stir extrusion (FSE) has gained particular interest as a promising method for producing wires from metal scraps, but recently, it was also employed for tube manufacturing. In literature, tube production via chip recycling often involves multi-step approaches, first consolidating/homogenizing the recycled chips and then extruding. In other cases, the tubes are manufactured directly from a bulk material, losing the sustainable goal. For this reason, this study aims to propose a single-step process in which aluminium chips are directly turned into a consolidated tube without any intermediate step. In addition, specific attention was given to the study of tool geometry, aiming to investigate the effect of a tapered tool's shape on the material flow and the overall process performance. Experimental tests were conducted to characterize the microstructure of extruded tubes and to calibrate numerical simulations employed for investigating process dynamics. Results revealed that the reduced contact diameter of the tapered tool generated lower processing temperatures but higher strain levels, fundamentally shifting the bonding mechanism from thermal assistance to mechanical dominance in oxide film breakage. Microstructural analysis demonstrated that the flat tool, characterized by predominant frictional heating and lower deformation, produced larger grain diameters due to thermally induced coarsening. Conversely, the tapered tool yielded significantly refined grain structures through severe plastic deformation and dynamic recrystallization under suppressed thermal conditions, indicating superior consolidation quality and enhanced particle bonding.

Introduction

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The industrial sector accounts for approximately one quarter of global CO2 emissions, reflecting the high energy intensity of large-scale material processing and transformation activities, with metals production representing a substantial contributor to this environmental burden due to its reliance on primary resource extraction and high-temperature processes [1]. Aluminum manufacturing, specifically, is responsible for roughly 3% of the world's direct industrial emissions [2]. As worldwide demand for aluminum continues to increase, driven by its favorable strength-to-weight ratio, corrosion resistance, and recyclability, which underpin extensive applications in the transportation, construction, and packaging sectors, the need to decouple resource consumption from economic growth has become increasingly critical to meet climate and sustainability targets [3]. Circular economy strategies provide a viable pathway to address this challenge by converting end-of-life products and manufacturing scrap into valuable secondary resources, thereby reducing dependence on energy-intensive primary production and limiting waste generation. Among circular economy approaches, recycling remains the dominant strategy for metallic materials due to their intrinsic ability to retain properties across multiple life cycles. Conventional aluminum recycling via remelting can reduce energy consumption by up to 90% relative to primary production, primarily because it eliminates the electrochemical reduction step and significantly lowers thermal energy requirements [4]. However, this route is affected by inherent inefficiencies, notably permanent material losses
arising from oxidation, dross formation, and alloy contamination during melting. These losses not only reduce material yield but also constrain the sustainability potential of conventional recycling, particularly for fine chips and low-quality scrap streams.
To address these limitations, solid-state recycling (SSR) techniques have emerged as innovative alternatives that completely bypass remelting, directly converting aluminum chips into semi-finished or finished products. SSR processes rely on the combined effects of temperature, pressure, and plastic deformation to achieve solid-state bonding between particles. The underlying mechanism involves the mechanical disruption of the stable oxide films naturally present on aluminum surfaces, enabling direct metal-to-metal contact and subsequent diffusion-driven bonding. Several SSR methodologies have been developed, each based on distinct mechanical principles to promote consolidation. Extrusion-based approaches, including direct hot extrusion and Equal Channel Angular Pressing (ECAP), impose severe plastic deformation that fragments oxide layers, enhances interparticle bonding, and refines the microstructure [5]. Friction-based methods, such as Friction Stir Extrusion (FSE) [6] and Friction Stir Consolidation (FSC) [7], exploit frictional heat generation and intense shear deformation to plasticize aluminum chips and weld them together under solid-state conditions. Sintering-based techniques, including Spark Plasma Sintering (SPS) and other field-assisted sintering methods, combine pulsed electrical currents, thermal energy, and compressive loads to promote diffusion bonding; however, they often require subsequent forming operations to achieve full densification due to limited strain introduction [8]. Among these approaches, direct extrusion processes represent a particularly promising SSR route, as they enable the transformation of loose chips into functional profiles, including hollow sections such as tubes, in a single continuous operation. The process mechanics involve feeding chips into a chamber where they undergo severe plastic deformation while being forced through an extrusion die, resulting in oxide dispersion, material densification, and effective solid-state bonding. The efficiency and quality of this transformation are strongly influenced by tool geometry, which governs material flow patterns, strain distribution, temperature evolution, and ultimately the resulting microstructural features and mechanical performance of the extruded product.
Tool design in extrusion processes has long been recognized as a critical factor affecting product quality in conventional manufacturing. However, its role in SSR processes, where the starting material consists of discrete particles with oxide-covered surfaces rather than continuous billets, introduces additional complexities. The tool geometry must not only guide material flow and shape the final profile but also ensure sufficient deformation to break oxide films and promote inter-particle bonding throughout the extruded section. Despite the importance of tool design in SSR extrusion, systematic investigations comparing different tool geometries and their effects on material consolidation and product quality remain limited in the literature.
The present study aims to elucidate the influence of tool geometry on the Friction Stir Extrusion (FSE) of seamless tubes produced from aluminum alloy chips, proposing a single-step manufacturing process able to convert chips directly into tubes. A comparative investigation is carried out on two distinct tool configurations: a conventional flat tool and a tool featuring a 20∘ chamfer at the shoulderdie interface. The flat geometry induces abrupt changes in material flow direction, whereas the tapered configuration promotes smoother material transition into the extrusion channel, altering local strain paths, pressure distribution, and flow continuity. The investigation integrates experimental campaigns with validated numerical simulations to provide a comprehensive assessment of both tool designs. Key process variables, including temperature distribution, strain accumulation, material flow patterns, and microstructural evolution, are analyzed to establish correlations between tool geometry and extruded tubes. Lastly, a microstructural analysis was carried out to enhance the effect of the different geometries. Understanding how these geometrical variations affect process mechanics, chip consolidation efficiency, and structural integrity of extruded tubes is essential for optimization and industrial scalability of SSR-based extrusion technologies.

Materials and Methods

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Experimental campaign development.

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

A dry milling operation was carried out on an AA2024-O aluminum bar ( 62 HV ) to produce machining chips, which were then cleaned in acetone for 30 minutes to ensure the best bonding performance and absence of contaminants. Successively, the chips were extruded using an ESAB LEGIO machine adapted for FSE (Figure 1a). A mass of 40 g of chips was loaded into a custom die chamber ( 76 mm in height and 25 mm in diameter), and two different tools were employed to consolidate the material into consolidated tubes (Figure 1b). Both tools had a nominal diameter of 23 mm ; however, one featured a 20∘ taper, reducing the effective diameter to 22 mm (Figure 1c). Two experimental campaigns were conducted at a rotational speed of 1000 rpm and a vertical load of 15 kN to evaluate the effect of the tapered design on the material properties and process operating. Lastly, a K-type thermocouple was positioned at one-third of the die height, 1 mm from the inner die wall, to monitor temperature evolution during processing and to support calibration of the numerical simulations. Microstructural characterization was performed on longitudinal sections. Samples were mounted, ground, polished, and etched using Keller's reagent ( 2 mLHF,3 mLHCl,5 mLHNO3,190mLH2O ). Observations were carried out with a GX51 Olympus optical microscope equipped with a 40× objective, and grain size was quantified using the mean linear intercept method (ASTM E11296) in ImageJ.

[picture]

Fig. 1. (a) Experimental setup, (b) process components of the FSE technique, and (c) tools adopted for extruding tubes.

Fig. 1. (a) Experimental setup, (b) process components of the FSE technique, and (c) tools adopted for extruding tubes.

Numerical campaign modelling.

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

The numerical campaign was conducted in the SFTC DEFORM 3D environment (Figure 2a). The die, backing plate, and tools were modeled as H-13 rigid bodies with mesh sizes of 30000,30000 , and 60000 elements, respectively. A shear factor of 0.2 and interface heat transfer coefficients (IHTC) of 11, 45 , and 45 W/mm2/K were assigned to model thermal exchange at the billet-tool, billet-die, and billet-backing plate interfaces. These values were determined through a coupled experimentalnumerical temperature and torque calibration process (Figure 3), in which the simulated temperature and torque profiles were iteratively compared with the measured thermocouple data until an acceptable agreement was achieved. The corresponding thermal and shear parameters were then adopted for all simulations. This calibration step is essential to ensure a realistic representation of the process. During calibration, variations in the IHTC led to only minor changes in the predicted temperature fields, whereas variations of the friction parameters had a significantly greater influence on both temperature evolution and material response. Given the computational complexity associated with the explicit modelling of individual chips, the chip mass was represented as a single porous billet (Figure 2b) following the Shima-Oyane constitutive formulation [9]. This formulation is selectable in the Deform software, and it is based on a corrected Von Mises criterion that takes into account the
presence of microvoids in the material matrix (by implementing the relative density " R " in the formulation). The initial relative density of the porous material was experimentally determined considering the geometries of the system. Specifically, it was calculated as the ratio between the mass of chips loaded into the die chamber and the corresponding internal volume. The resulting relative density was set to 0.7 , which represents the 70% of the density of the considered aluminium material. Lastly, to meet a compromise between computational time and quality of the results, the billet mesh consisted of 40000 elements, with a local refinement zone placed near the tool-material contact region (Figure 2b).

[picture]

Fig. 2. (a) Numerical simulation modeling of the experimental FSE setup and (b) example of the mesh window for the early-extruded numerical tube.

Fig. 2. (a) Numerical simulation modeling of the experimental FSE setup and (b) example of the mesh window for the early-extruded numerical tube.

[picture]

Fig. 3. Numerical calibration by means of experimental temperature and torque target profiles.

Fig. 3. Numerical calibration by means of experimental temperature and torque target profiles.

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

This section presents a comprehensive analysis of temperature evolution, strain accumulation, material displacement, and velocity field distributions. The numerical predictions are complemented by microstructural characterization of the tubes produced with both tool configurations, providing experimental validation and deeper insight into the correlation between process mechanics and material consolidation quality.
The numerical results obtained under steady-state conditions reveal significant differences in the thermo-mechanical behavior induced by the two tool geometries. As illustrated in Figure 4, the temperature and equivalent strain distributions along the longitudinal cross-sections of the extruded tubes demonstrate that the contact surface geometry significantly influences thermal generation. The reduced diameter of the tapered tool (Figure 4a) results in lower and more uniform frictional heating, yielding temperatures approximately 70∘C lower than those observed with the flat tool. Additionally, the larger contact area of the flat tool (Figure 4a) produces a distinct heat diffusion pattern in the vicinity of the tool-die clearance region (starting of the extrusion channel). Localized hightemperature zones are evident near the tool corners, which can be attributed to the severe material deformation imposed by the extrusion ratio and the associated flow restrictions within this region.

Strain distribution.

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Regarding strain distribution (Figure 4b), both tool geometries generate substantial equivalent plastic strain throughout the extruded material, which is essential for oxide film fragmentation and interparticle bonding. However, the flat tool induces higher strain concentrations near the tool-material interface and in the transition zone approaching the extrusion channel (in the range of 438-450). On the other hand, the tapered geometry promotes a more gradual strain accumulation, distributing deformation more uniformly across the material volume (in the range of 440-480). This difference in strain distribution patterns directly correlates with the material flow characteristics imposed by each tool configuration and has significant implications for consolidation quality and final tube properties.

[picture]

Fig. 4. (a) Temperature and (c) strain plot for the tube produced by the flat tool and (b) Temperature and (d) strain plot for the tube produced by the tapered tool.

Fig. 4. (a) Temperature and (c) strain plot for the tube produced by the flat tool and (b) Temperature and (d) strain plot for the tube produced by the tapered tool.

Temperature distribution.

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

The temperature distribution results indicate that the larger contact surface area associated with the flat tool geometry generates greater heating. This thermal condition induces two competing mechanisms: material softening due to elevated temperature and resistance to material flow. The tapered tool enhances material flow through its inclined surface but generates higher strain than the flat configuration. This is due to its lower peak temperature ( ~300∘C ) (Figure 4b), which reduces thermal softening and therefore requires greater plastic deformation to ensure proper consolidation (Figure 4d). Conversely, the elevated temperatures ( ~370∘C ) generated by the flat tool enhance material softening, resulting in reduced strain accumulation during processing (Figure 4c). Furthermore, strain localization patterns differ significantly between the two geometries. The flat tool exhibits concentrated high-strain regions at the corners ( ~445 ), adjacent to the extrusion channel, whereas the tapered tool generates a more distributed zone of elevated strain in the corresponding area. This distinct deformation behavior directly reflects the influence of the local temperature field on material flow characteristics.

Displacement and velocity plots.

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Material displacement and velocity fields along the extrusion direction (Z-direction) are presented in Figure 5, on the sectioned tube. The flat tool configuration (Figure 5a) induced significant flow obstruction, substantially affecting the displacement distribution in the tool vicinity. A wide region, extending several millimeters both above and below the tool, exhibited predominantly downward material displacement. This flow pattern indicates that the flat geometry redirected material away from the extrusion direction. In contrast, the tapered tool (Figure 5b) facilitated more streamlined material flow, resulting in visibly higher displacement along the extrusion direction and reduced downward-pushing effect. Moreover, the tube produced with the tapered tool exhibited a slightly shorter length compared to that obtained with the flat tool. This apparent contradiction can be justified by examining the strain distribution presented in Figure 4. Under the lower temperature conditions characteristic of the tapered tool configuration, greater plastic deformation is required to drive material through the extrusion die. At this point, the same volume of material undergoes more intensive straining, resulting in increased lateral flow and radial expansion rather than purely axial extension. This mechanism explains both the higher accumulated strain levels and the reduced final tube length observed in the tapered-tool case.

[picture]

Fig. 5. Vector plots of the total (a) displacement and (b) velocity along the extrusion direction for the flat tool and vector plots of the total (b) displacement and (d) velocity for the tapered tool.

Fig. 5. Vector plots of the total (a) displacement and (b) velocity along the extrusion direction for the flat tool and vector plots of the total (b) displacement and (d) velocity for the tapered tool.

The enhanced displacement observed with the tapered tool is directly attributable to the velocity distribution. As shown in Figure 5d, the extrusion rate is markedly higher in the tapered configuration, with velocity vectors concentrated along the tube wall and exhibiting substantial magnitudes. Material is efficiently channeled through the extrusion die, promoting continuous flow with minimal hindrance. Conversely, the flat tool (Figure 5c) generates lower extrusion velocities due to its larger contact surface, which impedes material flow. The velocity distribution in this configuration appears more uniform across the tube wall, though with reduced overall magnitude, reflecting the restricted material displacement previously discussed.
Experimental results confirm this trend. Under the same process parameters, the tube length was 50 mm with the flat tool and 56 mm with the tapered tool. The usable length (without the unbonded chips at the top and the unextruded material at the bottom) was almost 30 mm and 35 mm , respectively. The average thickness was approximately 1 mm for both tubes, with minor variations attributable to eccentricity from manual tool-die alignment.

Microstructure analysis.

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To assess the influence of tool geometry on microstructural evolution, transverse cross-sections were extracted from extruded tubes (Figure 6), and grain size measurements were performed close to the internal surface of the extruded tube. The analysis focused on regions experiencing distinct strain and temperature histories to elucidate the relationship between process variables and resulting microstructure.

[picture]

Fig. 6. Illustration of the longitudinal sections of the extruded tube, with the corresponding microstructural observations for the (a) tapered tool and (b) flat tool.

Fig. 6. Illustration of the longitudinal sections of the extruded tube, with the corresponding microstructural observations for the (a) tapered tool and (b) flat tool.

In the flat tool case (Figure 6b), characterized by higher processing temperatures and lower effective strain levels, frictional heating predominated as the primary heat source. Conversely, the reduced contact diameter of the tapered tool configuration (Figure 6a) minimized frictional heating. These distinct thermomechanical conditions directly influenced the observed microstructures. Material processed with the flat tool exhibited larger average grain diameters (avg. 4.3±0.5μ m ), attributable
to grain coarsening effects promoted by the elevated temperatures. The extended thermal exposure at higher temperature levels provided sufficient grain growth, partially offsetting any refinement achieved through plastic deformation. The lower strain accumulation in this configuration proved insufficient to counteract thermally induced coarsening, resulting in a relatively coarse final microstructure. In contrast, material extruded through the tapered tool displayed refined grain structure, with substantially smaller average grain diameters (avg. 4.1±0.3μ m ). The lower processing temperatures suppressed grain growth kinetics, while the intensified plastic deformation imposed severe shear strains that fragmented existing grains through dynamic recrystallization. The combination of reduced thermal coarsening and enhanced mechanical refinement resulted in a fine, equiaxed microstructure indicative of complete dynamic recrystallization.

Conclusion

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

In this paper, a single-step FSE process was successfully applied for producing consolidated tubes starting from aluminium chips. Two tools' geometries were tested through numerical simulations, along with a microstructural investigation. The following considerations were drawn:

A single-step FSE process is capable of producing fully consolidated tubes.

The numerical investigation revealed that the extended flat surface of the tool increases the frictional heat, promoting better material softening, but

The material flow through the clearance between the tool and the die is hindered more for the flat tool than the tapered one. Despite the tapered tool offering fewer obstacles to flow, the reduced heat resulted in a higher strain level, demonstrating the predominance of a mechanical-drive process over the temperature-driven one that characterized the flat tool geometry.

The reduced strain and high temperature combination, for the flat tool geometry, led to a visible grain coarsening compared to the tapered configuration, which limited the recrystallization phenomenon. A value of around 4.3 and 4.1μ m was observed for the flat and tapered tool, respectively.

Acknowledgment

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

This study was carried out within the MICS (Made in Italy - Circular and Sustainable) Extended Partnership and received funding from the European Union Next-Generation EU (PIANO NAZIONALE DI RIPRESA E RESILIENZA (PNRR) - MISSIONE 4 COMPONENTE 2, INVESTIMENTO 1.3 - D.D. 1551.11-10-2022, PE00000004). This manuscript reflects only the authors' views and opinions, neither the European Union nor the European Commission can be considered responsible for them.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.89.pdf

Bashmakov IA, Nilsson LJ, Acquaye A, Bataille C, Cullen JM, De La Rue Du Can S, et al., Contribution of Working Group III to the Sixth Assessment Report of the Intergovernmental Panel on Climate Change, (2022). https://doi.org/10.1017/9781009157926.013.

Aluminium - IEA, information on (accessed December 15, 2025).

Watari T, Nansai K, Giurco D, Nakajima K, McLellan B, Helbig C., Global Metal Use Targets in Line with Climate Goals, Environ Sci Technol (2020) 54:12476-83. .

Ashby MF. The materials life cycle, Materials and the Environment (2021) 41-64. .

Haase M, Ben Khalifa N, Tekkaya AE, Misiolek WZ. Improving mechanical properties of chipbased aluminum extrudates by integrated extrusion and equal channel angular pressing (iECAP), Materials Science and Engineering: A (2012) 539:194-204. .

Tang W, Reynolds AP. Production of wire via friction extrusion of aluminum alloy machining chips. J Mater Process Technol (2010) 210:2231-7. .

Li X, Baffari D, Reynolds AP. Friction stir consolidation of aluminum machining chips, The International Journal of Advanced Manufacturing Technology (2017) 94:5 2017;94:2031-42. .

Paraskevas D, Dadbakhsh S, Vleugels J, Vanmeensel K, Dewulf W, Duflou JR. Solid state recycling of pure Mg and AZ31 Mg machining chips via spark plasma sintering, Mater Des (2016) 109:520-9. .

Shima S, Oyane M. Plasticity theory for porous metals, Int J Mech Sci (1976) 18:285-91. .

This ePub file was created using machine-generated methods and automated software to convert the source text into an ePub format and may contain formatting inconsistencies from the original version. Formatting may vary depending on the device used.

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 99-109

doi: 10.4028/p-61pDv7

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-13

Accepted:
2026-02-24

Online:
2026-04-15

Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

The original version of this paper is available on https://www.scientific.net/MSF.1185.99.pdf

Clark Austin 1,4,a, Barnett Philip 2, b, Mamros Elizabeth 3,c, Young Brian 4, d and Ragai Ihab 4,e*

1 GROB Systems Inc., Bluffton, OH 45817, USA

2 University of Tennessee, Knoxville, TN 37996, USA

3 Bucknell University, Lewisburg, PA 17837, USA

4 The Pennsylvania State University, Erie, PA 16563, USA

a  aclark@grobsystems.com, b  pbarnet3@utk.edu, c  elizabeth.mamros@bucknell.edu, dbay101@psu.edu, e*ihab.ragai@psu.edu

*corresponding author

Keywords: FSW, carbon fiber reinforced polymers, polyphenylene sulfide, microstructure.

Abstract

Friction stir welding (FSW) is widely used to bond metals and polymers to create large structures from standardized geometric components. However, FSW has found more limited use in traditional continuous fiber reinforced composites due to the risk of damage to the reinforcing fibers, including fiber misorientation and fragmentation. The process is more amenable to discontinuous fiber composites, which often take the form of injection molding compounds reinforced with short, milled fibers, for which the mechanical performance is significantly degraded. An emerging class of materials, recycled carbon fiber nonwovens, contain long discontinuous fibers that imbue their composites with mechanical performance that bridges the gap between injection molding compounds and continuous fiber laminates, while also decreasing the energy footprint of the materials. This work evaluates this class of materials, alongside injection molding compounds, to develop new insights into the bonding of composite structures using friction stir welding. The bridging effect of discontinuous fibers across the bondline is evaluated using optical microscopy to link processing conditions and fiber length to the resulting performance. Fiber migration was observed in the weld area, though mechanical interlocking was the primary mechanism of bonding in the weld zone. While samples failed at a fraction of the neat materials nominal strength, increased fiber length was found to have a beneficial effect on the apparent tensile and lap shear strength on welds in this study. The new insights gained represent an important step towards the adoption of the FSW process as a means of rapidly manufacturing large composite structures made of carbon fiber/PPS or other fiberreinforced polymers to enable deployable structures from standardized geometric components.

Introduction

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Fiber-Reinforced Polymers.

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In polymers with fiber reinforcement, the bridging effect is observed when fibers cross from one face of a delamination to the other, beneficially increasing the strength of the polymer [1]. Both fiber length and orientation have been found to play a role in the strength of reinforced polymers [2]. In particular, fiber length plays a major role in the strength of composites, with fibers below the critical shear transfer length failing due to transfer stress from the matrix to fiber. Below this critical length, fiber pull-out is the dominant failure mechanism, while above this length, fiber fracture dominates. For an ideal composite with mono-strength fibers and perfect dispersion, the strength will theoretically be independent of fiber length beyond the critical fiber length, but such materials do not exist in practice. Instead, real composite materials exhibit a distribution of fiber lengths and fiber strengths, as well as inhomogeneous dispersion leading to stress concentrations that reduce strength. For example, in recent work by Barnett et al., carbon fiber reinforced polyphenylene sulfide (CF/PPS)
composite tensile strength and modulus were shown to exhibit strong fiber length dependence [3]. For injection molding (IM) compounds, they demonstrated that doubling the fiber length resulted in a 28% increase in tensile strength despite a substantial reduction in total achieved fiber volume fraction ( 36% for the shorter fibers and 12% for the longer).

Increasing the fiber length far beyond the critical fiber length yields much higher performance at similar fiber volume fractions, as shown in Table 1. Recent efforts have focused on the development of long discontinuous fiber composites that offer greater formability for complex geometries than continuous fiber laminates while retaining moderate performance that exceeds IM compounds. Examples include platelet composites [4] and quasi-isotropic nonwovens [5], which offer unique manufacturing and performance benefits, making them well positioned for low-cost, high-volume production. Compared to platelet composites, which use precisely chopped virgin fiber prepreg, nonwovens are amenable to a range of fiber lengths and feedstocks, including recycled fiber, and achieve similar tensile strength at a significantly reduced fiber volume fraction, further reducing cost.

Table 1. Performance of common CF/PPS composite formats.

Material | Format | Fiber Volume Fraction [%] | Tensile Modulus [GPa] | Tensile Strength [MPa]
Toray Cetex TC1100 [6] | Woven Fabric | 50 | 58 | 752
Mitsubishi Techtron PSBG PPS [7] | Injection Molding Compound | 38 | 6.3 | 35
Toray Cetex MC1100 [8] | Platelet Molding Compound | 59 | 41.4 | 207
Carded Organosheets [5] | Nonwoven | 29 | 25 | 204

Friction Stir Welding.

The original version of this paper is available on https://www.scientific.net/MSF.1185.99.pdf

When joining materials by FSW, the coefficient of friction (CoF) of the workpieces must be considered, as the heat generated between the workpiece and the rotating tool shoulder is the primary source of heat input for the weld [9]. The heat generated from the rotation of the tool shoulder, by way of sliding and sticking forces, softens the workpiece material allowing for plastic deformation along the rotational path of the tool [10]. The CoF of CF/PPS is typically lower than neat PPS [11], which reduces the thermal effect from the rotating tool shoulder. Especially in thermoplastics, the CoF and thermal conductivity are lower than metals, which drives the requirement for pre-heating in order to reach temperatures suitable for plastic deformation [12][13]. It is for these reasons that parameter selection and pre-heating are critical factors in creating a sound weld in CF/PPS. When FSW polymers, the weld bond relies on diffusion across the polymer-polymer interface in intimate contact above the glass transition temperature, Tg[14]. The reptation time of the polymer, an inversely related function of temperature, must be reached to achieve complete bonding across the interface. In CF reinforced polymers, fiber characteristics such as length, orientation, and fiber volume fraction are also very influential on joint strength [15]. It is noted that welded CF polymers often exhibit lower strength than their virgin counterparts due to disruption and a transverse flow of fibers in relation to the seam. However, if fibers adequately span the weld line and sufficient polymer diffusion occurs, the strength of CF reinforced welds should improve.

The work presented herein evaluates the feasibility of bonding composite structures using FSW. The findings from this study will be used to further develop joining practices of CF/PPS for use in larger composite structures than can be made by molding processes. Limitations on mold size and a knowledge gap in joining CF/PPS by FSW, especially that of the recycled nonwoven discontinuous type, hinders the ability for the manufacturing of large robust structures and components. By developing a feasible and practical method for joining CF/PPS by FSW, larger lightweight components can be assembled at lower cost for use in fields, such as aerospace and defense, where strength-to-weight ratios prove critical.

Materials and Methodology

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Materials.

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To fully understand the relationship between fiber length and bridging at the bondline, multiple PPS formats were tested in similar and dissimilar lap joints. Injection molding was used to manufacture CF/PPS (Electrafil PPS CF50 3DP, TechmerPM) plates containing milled CF ~135μ m in length (herein referred to as IM) along with neat PPS plates (Fortron 0203, Celanese). Square plaques ( 152.4 mm×152.4 mm ) were molded in an Arburg 570A 2000-800. The front barrel temperatures were set to 316∘C. The mold was maintained at 150∘C with the use of a Regloplas P 160 M steam temperature controller. A fill rate of 100cc/sec was used to fill the cavity. The natural PPS required 14.12 MPa and the carbon fiber filled PPS required 118.84 MPa to fill the cavity 95% full. Prior to molding the material, the pellets were dried at 65∘C for a minimum of 8 hours. The plaques were molded at a 44second cycle.

Additionally, compression molding was used to manufacture CF/PPS plates from recycled "short" chopped carbon fiber nonwovens (Gen 2 Carbon) with lengths of ~38.8±22.2 mm (herein referred to as CMS). Continuous unidirectional carbon fiber prepregs (P20053-12X, Teijin Carbon) were used to manufacture quasi-isotropic laminates (herein referred to as CML). Compression molding was done at 300∘C for 5 minutes and then cooled under 3 MPa pressure for the recycled fiber nonwovens and following the manufacturer's recommended cycle for the prepregs. Materials tested along with experimental joints are shown in Table 2.

Table 2. Samples used in experiments.

Sample # | Top Plate | Bottom Plate
1 | IM | IM
2 | IM | CMS
3 | IM | CML
4 | CMS | CMS
5 | CMS | IM
6 | CML | CML
7 | CMS | Neat

Tensile Testing.

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Tensile testing of the CML material was completed in accordance with ASTM D3039M standard using straight-sided specimens [16]. The strain was measured using 2-D digital image correlation (VIC 2D, Correlated Solutions). The tensile properties of the CMS materials were acquired as part of previous work [5]. The goal of this test was to evaluate the baseline performance of the different materials for comparison. Additionally, tensile testing samples were machined out the IM plaques in accordance with ASTM D638 [17]. The tests were conducted using an Instron Universal Testing Machine with a 30 kN load cell. The tests were run at 1.0 mm/min until failure. Strain was calculated using the crosshead displacement. Three samples per plaque per material were tested.

IZOD Impact Testing.

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The impact performance was completed according to the ASTM D256 standard [19]. The goal of this test was to elucidate the effect of fiber length on the impact performance as a means of quantifying fiber bridging within each material. An impact energy of 11.3 J was used to test the CMS and CML samples (IT503, Tinius Olsen). An impact-type hammer with a pendulum weight of 1.0 kg [2.2 lb] and energy at impact of 6.1 J [4.5 ft-lb] was utilized for IM samples. Results are reported in terms of energy absorbed per unit of specimen cross-sectional area under the notch. Three samples from each IM material were tested.

Friction Stir Welding.

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A lap joint configuration was used for all tests. Parameters used for FSW were a feed rate of 5 mm/min, tool rotation speed of 1500 rpm in a counterclockwise direction, a shoulder plunge depth 1
mm below material surface and a 5 -second dwell before welding. All welding was performed in a horizontal 3-axis CNC machine. A weld length of 100 mm was utilized for all samples. The FSW tool used had a pin length of 5 mm , pin diameter of 6 mm , shoulder diameter of 20 mm and a threaded triangular truncated fluted pin profile. All plates were 3.175 mm in thickness, creating a weld stack up of 6.35 mm ; thus, giving 95% pin penetration in the weld. A schematic of the pin depth in the workpieces and diagram of the FSW tool is shown in Fig. 1. A Creality Ender-3 S1 3D printer was modified for use as a preheating unit to achieve 85∘C in the samples before welding. Preheating was accomplished on the bottom sample via the print bed under an aluminum plate while the top sample utilized the heating element intended for the printer nozzle. A steel bar was clamped between the top sample and nozzle heating element to aid in thermal distribution. An image of the experimental setup is shown in Fig. 2. Both heating elements were set to 100∘C and monitored until the material reached 85∘C. An Optris 640i LT infrared (IR) camera with a wavelength range of 8−14μ m was used to monitor and collect the temperature distribution during welding as shown in Fig. 3. Emissivity was set to 1.0 for all samples as the top material surface was black in color. An unobstructed line of sight between the IR camera and workpieces was set to 1 m . Tool occlusion was noted as a factor in the experimental setup. The maximum measured temperatures were likely lower than the actual temperature under the tool during welding. While this is a limitation of IR measurement, trends in the measured temperature values among samples were compared equivalently. An AMTI MC12-6-4000 transducer was used to collect force readings in X,Y, and Z directions during welding.

[picture]

Fig. 1. FSW tool pin depth diagram.

Fig. 1. FSW tool pin depth diagram.

[picture]

Fig. 2. Schematic of experimental setup.

Fig. 2. Schematic of experimental setup.

[picture]

Fig. 3. Thermal image during welding of sample 6 with maximum temperature value recorded at that instant.

Fig. 3. Thermal image during welding of sample 6 with maximum temperature value recorded at that instant.

Lap Shear Testing.

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Lap shear testing of welded specimens was completed according to the ASTM D5868 standard [18] using specimens with a modified geometry (nominally 76 mm total length with 25 mm×25 mm overlap) and G-10 fiberglass doublers. The samples were gripped 20 mm from each end and loaded at a rate of 13 mm/min. A minimum of three specimens were tested for each sample type with the exception of sample 1 for which only a single specimen was available.

Optical Microscopy.

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Optical microscopy was used to evaluate the surface roughness and microstructures of the composites (VHX-7000, Keyence). The surface roughness of the weld area was analyzed using the optical profilometry feature of the microscope to report the arithmetical mean height ( Sa ), the maximum height (Sz), and the root mean square height (Sq). A minimum of four surfaces were measured for each sample type. Samples were also extracted from the end of the weld region, mounted in epoxy, and polished to observe the weld microstructure. The weld area size in the base material was measured using the open-source Fiji image analysis software [20]. The weld area size was measured for one specimen per sample type.

Differential Scanning Calorimetry (DSC).

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Differential scanning calorimetry was used to study the crystallization and melting behavior of the composites. Samples of 4 to 7 mg mass were heated from 20 to 350∘C in a nitrogen atmosphere at 20∘C per minute using a Perkin Elmer DSC 8500. The samples underwent two heating cycles. The enthalpy of melting was taken from the first heating cycle to identify the crystalline percentage using equation (1):

Xc=ΔHmΔHf(1−Wf)(1)

where ΔHm is the measured melting enthalpy, ΔHf is the heat of fusion of PPS taken as 88.37 J/g, and Wf is the weight fraction of fiber in the sample. The melting onset and peak temperatures were identified from the second heating ramp, which provides more consistent data due to enhanced thermal contact but loses information pertaining to the thermal history of the as-processed material. A minimum of three samples each were extracted from within the weld zone and on the edges furthest from the weld zone. The goal of this test was to identify if the FSW process led to changes in the thermal properties that might indicate material degradation due to heat and shearing in the weld zone.

Results and Discussion

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Baseline Properties.

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The tensile and IZOD impact properties of the baseline materials are provided in Table 3. As anticipated, the highest performing material was the CML, whose continuous fibers more effectively translate loads across the composites. The modulus of the quasi-isotropic CML laminates was quite similar to that of the CMS, due to the mildly anisotropic nature of the nonwovens (anisotropy ratio of roughly 2:1 [5], yielding a quasi-isotropic laminate modulus of approximately 25 GPa ). The strength of the continuous fiber laminates was much improved, largely attributed to the reduced stress concentrations that cause premature failure in discontinuous fiber composites (as evident from the reduced failure strain). Still, for stiffness-driven applications, CMS laminates may be desirable, especially in cases such as thin skins bonded to CML stiffeners.

For the IM samples, tensile results are also shown in Table 3. The addition of carbon fibers to the PPS matrix reduced the elongation and ultimate strength. This could be attributed to the increased stiffness and brittleness of the material. Additionally, the decline could also be attributed to stress concentration at the ends of the fibers and agglomeration within the matrix. The IZOD results indicate that milled carbon fibers enhance the toughness of PPS. This could be attributed to the carbon fibers acting as obstacles to the propagation of cracks within the PPS matrix. When a crack reaches a fiber, it gets deflected along the fiber-matrix interface and causes the fiber to debond or causes the fiber to break. Each of these scenarios consumes a notable amount of energy, thereby increasing its overall toughness, even while becoming more brittle in a simple tensile test.

Table 3. Mechanical Performance Results for PPS.

Material | Fiber Volume Fraction [%] | Tensile Strength [MPa] | Tensile Modulus [GPa] | Failure Strain [%] | IZOD Impact Resistance [kJ/m²]
Neat | 0 | 62.4 ± 10.0 | 2.3 ± 0.1 | 3.19 ± 0.62 | 1.55 ± 0.3
IM | 36 | 49.7 ± 2.1 | 6.2 ± 0.9 | 0.98 ± 0.22 | 3.11 ± 0.4
CMS | 32 | 268 ± 32.3 | 34 ± 2.6 | 0.80 ± 0.09 | 15.3 ± 1.0
CML | 53 | 594.8 ± 7.1 | 38.8 ± 0.9 | 1.61 ± 0.07 | 99.4 ± 7.9

Thermomechanical Analysis.

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Thermal and longitudinal force measurements were collected for all samples during welding. Maximum temperature values are shown in Table 4. Force data in the longitudinal direction, that is in the direction of the weld seam, is shown in Table 5. Force values are partitioned by first quartile value (Q1), median, third quartile value (Q3) and maximum along with standard deviation. In general, when IM samples were welded in the top plate position (samples 1-3), lower temperatures were generated than with CM samples in the top plate position (samples 4-7). This can be attributed to lower frictional forces as shorter fibers slide past each other in the weld zone. The decrease in temperature for sample 7 (CMS/Neat) is attributed to reduced friction in the neat PPS portion of the weld zone. In general, higher variation in the weld force was observed for samples containing injection molded portions (samples 1-3, 5, and 7). The lowest variation in weld force was for sample 6, which also contains the most homogenous microstructure. The cause of these variations may be changes in material flow that vary spatially alongside variable microstructure.

Table 4. Temperature data ( ∘C ).

Sample | Maximum
1 | 214.9
2 | 229.9
3 | 215.2
4 | 275.0
5 | 275.0
6 | 275.0
7 | 260.3

Table 5. Longitudinal force data (N).

Sample | Q1 | Median | Q3 | Maximum | Mean | Std Dev (σ)
1 | 18.9 | 46.7 | 75.3 | 134.2 | 48.6 | 33.2
2 | 25.2 | 49.7 | 77.2 | 133.5 | 51.8 | 31.1
3 | 39.6 | 64.8 | 91.6 | 143.6 | 63.2 | 36.9
4 | 45.0 | 70.9 | 102.7 | 130.0 | 69.5 | 38.2
5 | 60.6 | 92.3 | 130.1 | 160.8 | 89.1 | 43.4
6 | 54.6 | 72.5 | 89.4 | 106.7 | 67.2 | 28.4
7 | 13.4 | 37.6 | 70.5 | 113.6 | 42.2 | 32.6

Lap Shear Testing.

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All welded samples failed in invalid failure modes - rather than failing in shear at the lap joint, the failure was in tension at the weakest side of the welded structure. In most cases, this was the weld side (samples 2, 3, 4, and 6); however, when a strong material was welded to a weaker base material (samples 5 and 7) the failure transitioned to the base side. Sample 1 was an exception to this result, failing instead by pull-out of the welded region and separation into the two original pieces, indicating that the milled carbon fiber was too short to effectively bridge the weld interface. The three failure modes are shown in Fig. 4. The observed behavior was consistent for all specimens within a sample set. In all cases, the weld appeared to be less of an adhesive bond and more of a mechanical interlocking mechanism. As such, the apparent tensile and shear strengths have been reported based on the failure side thickness and width and the overlap area, respectively. Apparent tensile and lap shear strength are graphed in Fig. 5. These bar graphs show the mean strength and standard deviation for each sample type. The mechanical strength of the welds in this study was far below the strength of the unwelded materials.

[picture]

Fig. 4. Failure modes of samples 1 (a), 2 (b), and 7 (c) showing failure is dominated by tension in the weakest side of the sample.

Fig. 4. Failure modes of samples 1 (a), 2 (b), and 7 (c) showing failure is dominated by tension in the weakest side of the sample.

[picture]

Fig. 5. Apparent (a) tensile and (b) lap shear strength graphs.

Fig. 5. Apparent (a) tensile and (b) lap shear strength graphs.

Differential Scanning Calorimetry (DSC).

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DSC results indicate a slight decrease in crystallization and slight increase in the melting onset and peak temperatures in the weld area for most of the samples, as shown in Fig. 6. The standard deviation is represented by black error bars. The results indicate thermo-oxidative degradation, in which the PPS begins to crosslink, thereby impeding crystallization and increasing transition temperatures [21]. This degradation is most prevalent at temperatures well above the melting temperature of PPS in the presence of oxygen. While the temperature data acquired during FSW testing does not show such temperatures, the true temperature in the weld zone may be masked by the tool and surrounding sample. Given the welding occurred in standard atmosphere and optical microscopy results show evidence of melting, it is hypothesized that these small changes are attributed to thermo-oxidative degradation during welding.

[picture]

Fig. 6. DSC for welded & unwelded areas of each material.

Fig. 6. DSC for welded & unwelded areas of each material.

Optical Microscopy (OM).

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Transverse weld sections were observed by way of optical microscopy (OM) to analyze the interaction between fiber length and orientation in and around the weld zone. Fig. 7a shows sample 6, consisting of quasi-isotropic layers of continuous unidirectional carbon fiber, after welding. Outside of the weld zone, the laminate sequence is retained. Meanwhile, the stir zone (SZ) exhibits severe fiber disruption along with voids on both sides of the thermomechanically affected zone (TMAZ) and in the root of the weld. Sample 4 exhibited similar fiber disruption within the SZ as shown in Fig. 7b. Within the SZ, plastic deformation caused by the rotation and longitudinal motion of the tool likely damaged fibers as evident from lateral voids indicating fiber pull-out. It was observed that both fiber length and orientation were severely disrupted by the FSW process, with post-weld mechanical conditions relying heavily on the degraded polymer properties. Fig. 7c shows the transverse weld section of sample 7. Evidence of fiber migration within the weld can be seen from the top plate into the bottom plate. While not as clearly apparent in the other samples, due to difficulty in tracking fiber migration, this indicates that the bridging effect likely contributes to the differences in apparent lap strength for composites with longer fibers. The cross-sectional area of the weld zone was mostly constant for all composite base materials (ranging from 4.9 to 6.1 mm2 ) but increased substantially when the base material did not contain reinforcing fibers ( 8.9 to 10.1 mm2 ), indicating that fiber migration is strongly dependent on the material surrounding the SZ.

[picture]

Fig. 7. Transverse section of samples 6 (a), 4 (b), and 7 (c) weld regions.

Fig. 7. Transverse section of samples 6 (a), 4 (b), and 7 (c) weld regions.

Surface Roughness.

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Optical profilometry was utilized to measure the surface roughness. The results are listed in Table 6 along with the weld area. In general, samples with CMS on the weld side had higher Sz than IM samples, while CML welds showed the lowest peak-to-valley variation. The Sa results were generally more similar between materials, with the exception of samples 5 and 7, which also exhibited the largest Sq values. This is likely due to lower stiffness in the bottom plates resulting in greater fiber migration.

Table 6. Optical microscopy results.

Sample # | Sa (μm) | Sz (μm) | Sq (μm) | Weld area (mm²)
1 | 135.6 ± 31.8 | 1157.8 ± 95.0 | 184.5 ± 25.0 | -
2 | 139.9 ± 44.0 | 1155.9 ± 288.9 | 177.3 ± 52.9 | 5.655
3 | 98.4 ± 14.6 | 859.9 ± 233.2 | 127.5 ± 25.0 | 4.919
4 | 108.6 ± 36.7 | 1070.1 ± 112.5 | 147.1 ± 38.0 | 4.917
5 | 172.5 ± 67.2 | 1665.2 ± 701.5 | 240.0 ± 102.9 | 4.589
6 | 117.7 ± 12.5 | 840.5 ± 142.6 | 147.8 ± 19.7 | 5.378
7 | 272.2 ± 63.1 | 2260.2 ± 744.8 | 353.9 ± 80.2 | 9.459

Conclusion

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In the current work, polyphenylene sulfide (PPS) with varying lengths of carbon fiber reinforcement were joined by FSW to measure the effect of fiber length on mechanical and microstructural properties. While fiber migration was observed by optical microscopy, the absence of characteristic shear failure and minimal fiber bridging does not inherently convey that fiber integrity was transferred through the bond line. There is evidence that continuous plastic deformation in the weld zone
diminished fiber length and orientation while the thermal input during welding caused polymer degradation, reducing mechanical strength within the weld zone. Although the apparent tensile and shear strength of the welded samples in this study was a fraction of the parent material, CF/PPS with longer fiber lengths did exhibit a positive relation to mechanical strength.
Key findings from this study include:

Welding generally increased melting onset and peak temperature, indicating thermal degradation of the polymer matrix.

The FSW process led to macrostructures dominated by mechanical interlocking rather than adhesive bonding.

Severe fiber disruption and misorientation were present in the weld area, regardless of fiber length.

A positive relationship between fiber length and weld strength was observed with fiber migration visible in the weld area, indicating some bridging effect.

Future Work

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Future work will continue to focus on optimizing FSW parameters in pursuit of mechanical properties more closely aligned with the CF/PPS parent material.

Acknowledgement

The original version of this paper is available on https://www.scientific.net/MSF.1185.99.pdf

The authors would like to thank Arame Cisse, Broc Dawson, Kristofer Laser Jr., Senudin Mehinovic, and Mason Sheeley along with Andrew and Christopher Bartlett of Penn State Behrend for their assistance with the experimental work. The authors would also like to thank Gen 2 Carbon for their donation of the recycled carbon fiber nonwoven used in this work. We acknowledge the partial support of the US National Science Foundation in the form of a Travel Support Grant [CMMI 2511050], which enabled the authors to attend and present at this conference.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.99.pdf

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 111-125

doi: 10.4028/p-1wZ1l6

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-16

Accepted:
2026-02-24

Online:
2026-04-15

Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

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MATHEW Amith Tom 1,a*, PIWEK Armin 1, b, MAJER Jakub 2,c, MOHNFELD Norman 1, d, BEHRENS Bernd-Arno 1,e and UHE Johanna 1,f

1 Leibniz University Hannover, Institute of Forming Technology and Machines, An der Universität 2, 30823 Garbsen, Germany

2 Poznań Institute of Technology, Łukasiewicz Research Network, Ewarysta Estkowskiego St. 6, 61-755 Poznań, Poland

a* mathew@ifum.uni-hannover.de, b piwek@ifum.uni-hannover.de, cjakub.majer@pit.lukasiewicz.gov.pl, dmohnfeld@ifum.uni-hannover.de, èbehrens@ifum.uni-hannover.de, f  uhe@ifum.uni-hannover.de (*corresponding author) 

Keywords: Rotary friction welding, FEM, Tailored forming

Abstract

Rotary friction welding (RFW) is a solid-state process used to join similar and dissimilar materials, such as steel and aluminium. In RFW, interface temperature development and its distribution are essential factors influencing material bonding. It governs bond strength, the formation of intermetallic phases (IMPs) and the evolution of the heat-affected zone (HAZ). Thus, precise prediction of temperature distribution is vital for the reliable design and optimization of the RFW process, as well as for the prediction and control of IMP formation. This work presents an experimental investigation of the thermo-mechanical behaviour of EN AW-6082 and 20 MnCr 5 during RFW in addition to a corresponding novel numerical modelling framework. A systematic parameter study was conducted to evaluate the influence of the friction pressure, friction time, forging pressure, forging time and rotational speed on the peak temperature, sledge path and flash formation. In-situ temperature measurements were performed using thermocouples (TC) embedded in the steel component, while axial force, displacement and rotational speed were recorded. The results demonstrate that, within the investigated parameter ranges, the rotational speed is the dominant factor governing frictional heat generation and the peak temperature, while the friction pressure primarily influences the sledge path. In parallel, a 2D axisymmetric finite-element model with a user-defined subroutine was developed to compute the heat flux based on process parameters and contact conditions, providing a transparent and extensible numerical framework for RFW. The experimental findings establish a robust basis for the calibration and future validation of the numerical model.

Introduction

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The automobile industry aims to reduce CO2 emissions and thereby improving the energy efficiency by focusing on lightweight automobiles [1,2]. However, every component cannot be replaced by lightweight materials like aluminium. Critical structural components require materials that can withstand high temperatures and mechanical loads [1,3]. Tailored forming techniques are being developed to produce semi-finished hybrid parts which can fulfil these requirements [ 3,4 ]. Automotive manufacturing commonly employs both steel and aluminium alloys, either as individual materials or in hybrid component combinations. In particular, the development of steel-aluminium hybrid components has gained increasing attention due to their ability to combine high strength with weight reduction. This interest has motivated extensive research into suitable joining technologies, especially solid-state processes such as friction welding [5]. They are particularly favoured due to their cost-effectiveness and favourable mechanical properties. RFW technique is a plausible method of material joining especially when it comes to dissimilar material combinations [3,4].

Welding processes can be classified into fusion welding and solid-state welding, based on whether the base material is melted or not. In solid-state welding, heat is generated in the joining zone by external mechanical loading, whereas in fusion welding the joining zone is heated by an external heat
source [6]. RFW is a solid-state joining technique widely applied in aerospace, automotive and energy sectors for producing high-integrity joints with limited distortion [7]. In RFW, heat is generated through the conversion of mechanical energy into thermal energy at the interface between the rotating and axially loaded components. One component is rotated under axial pressure against a stationary counterpart and the resulting frictional heat rapidly raises the interface temperature, leading to localised material softening [7].

The RFW process is generally divided into two main stages: the friction phase and the forging phase [3,7]. The rotational speed is initially maintained at a constant value during the friction phase, while a predefined friction pressure is applied to generate heat at the interface. At the end of the friction phase, the spindle rotation is abruptly stopped and the axial force is increased to the forging pressure. This forging phase promotes material consolidation and joint formation promoting metallurgical bonding. Flash formation around the joint region demonstrates that significant plastic deformation and material expulsion occurs during the forging phase [7]. Given its advantages over conventional methods, including reduced costs, shorter welding times, higher accuracy, excellent bond strength resulting from its fine-grained structure, fewer welding defects and suitability for joining dissimilar materials, this technique offers significant benefits [8].

The rotational speed, friction time, friction pressure and forging pressure are important parameters that can be used to control the secondary parameters like friction heat and material flow, influencing the joint strength and weld quality [3]. At elevated rotational speeds and higher friction pressures combined with extended friction times, the heat generated at the interface increases, leading to a more uniform temperature distribution across the joining zone. However, the central region typically exhibits lower temperatures due to its lower relative velocity, resulting in reduced frictional heat generation near the axis of rotation [9]. Excessive heating under these conditions can promote pronounced thermal softening and grain coarsening in the weld zone, which has been linked to a deterioration of mechanical properties, including reduced joint strength. In contrast, more moderate combinations of rotational speed, friction pressure and friction duration produce lower peak temperatures and avoid excessive softening, thereby establishing a more stable process window in which higher joint strengths have been observed [ 10,11 ]. The large differences in thermal conductivity, flow stress and plastic deformation behaviour between steel and aluminium make thermo-mechanical control difficult during RFW. These mismatches strongly influence heat generation, flash formation, HAZ evolution and critically, the IMP formation at the interface [8,12]. Although the aluminium alloy remains below its melting temperature throughout the RFW process, diffusion mechanisms are activated, leading to the development of a layer consisting of intermetallic phases [3]. Although thin IMP layers can contribute to metallurgical bonding, excessive growth leads to extreme brittleness, impaired ductility and reduced fatigue performance [3,13]. IMP thickness and morphology are strongly governed by the thermal cycle generated during the friction and the axial pressure during the forging stages of RFW. Predicting the temperature evolution during friction welding is crucial, as the thermal field directly controls material flow and microstructural changes. The accurate prediction and control of heat generation are thus central to predicting and minimising intermetallic formation and ensuring joint integrity [3,12,13]. Since the key thermal quantities cannot be measured directly with standard machine sensors, process simulation provides essential insight into the transient temperature distribution within the joining zone [10].

Thermo-mechanical numerical modelling is particularly valuable for RFW processes because the welding occurs over short durations, involves highly coupled thermal and mechanical phenomena and is sensitive to process parameters [8,14]. Finite-element (FE) simulations provide a practical and cost-effective approach to quantify the thermo-mechanical behaviour, including temperature distribution, plastic strain, contact pressures and heat generation at the interface that are difficult to measure experimentally. Consequently, such simulations facilitate a deeper understanding of the underlying process physics and allow systematic optimization of welding parameters to improve joint quality and mechanical performance [14,15].

Numerical modelling of RFW used prescribed radial heat fluxes and rotational-speed dependencies, which revealed steep thermal gradients near the interface [16]. However, these early thermal models lacked mechanical coupling and neglected friction and contact physics. Their inability to predict flash formation and residual stresses further constrained their use for dissimilar material combinations [14]. The seminal 2.5D work of Moal & Massoni introduced axisymmetric FE models that retained circumferential kinematics. These models used adaptive remeshing to enable the predictive simulation of the rotational deceleration and sledge path during the welding of similar material combinations. However, the approach did not consider post-weld cooling, or interfacial phase changes and the experimental calibration was limited [17]. Davé et al. developed analytical thermal models for dissimilar welds where energy input was taken from measured angular speed curves and used to compute thermal profiles [18]. Later, Wang et al. and Lee et al. measured angular/torque histories to define interface thermal boundaries for coupled FE models [11,18]. D'Alvise et al. implemented a coupled thermo-mechanical RFW module with a pressure-speedtemperature friction law and validated it for similar and dissimilar nickel-based alloys. Although the model reproduced general parameter sensitivities, it systematically over-predicted weld time, forging and interface temperatures compared with experiments across weld conditions. This highlights the challenges in accurately modelling interface behaviour, particularly for dissimilar materials with strong thermal-mechanical contrasts [19]. Despite these advances, most studies remain focused on similar metal systems. Bennett's integrated FE study that explicitly included phase transformations and validated residual stress predictions represents a mature template for modern coupled RFW simulations. However, the investigation was restricted to monomaterial welding conditions involving steel [14]. Recent work by the authors on C22.8-41Cr4 friction welds employed the inbuilt Simufact friction welding module to develop a substitute thermo-mechanical model. However, the predictive capability of this approach decreased noticeably when process parameters were varied, with significant deviations observed during validation. More critically, the model relies on a simplified scaling factor, (described as the Y-Factor) in the frictional heat generation and energy-partitioning formulation [8]. Currently, there are insufficient detailed experimental data for dissimilar material combinations, particularly steel-aluminium joints, to reliably calibrate and validate predictive thermomechanical models. Therefore, experiments are required to investigate temperature evolution, thermal gradients and heat generation behaviour in these joints, as well as to provide insights into material flow and their dependence on process parameters. The data obtained from these experiments will form the foundation for developing accurate coupled models capable of capturing intermetallic formation in dissimilar metal systems.

To address this limitation, a fully physics-based 2D axisymmetric FE model was developed using a user-defined subroutine, providing explicit access to the interface friction laws, heat generation and material response. The development, calibration and validation of the numerical model require reliable experimental input, including accurate material data and experimental measurements acquired under representative process conditions. Therefore, RFW experiments were conducted using 20 MnCr 5 steel and EN AW-6082 aluminium alloy workpieces under various process parameter combinations (PCs). In-situ temperature measurements were performed during the welding process to capture the thermal evolution near the joint interface. A systematic parameter study was carried out based on the resulting experimental dataset to analyse the influence of key process parameters on the thermal and mechanical response of the joint. In order to evaluate the developed subroutine, the numerically calculated temperature curves were compared with the experimental curves.

Methods - Material Data

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

Comprehensive material data for EN AW-6082 and 20 MnCr 5 , including temperature-dependent mechanical and thermal properties, have been previously investigated and reported in the prior studies from the same research framework [20]. The experimentally determined flow curves were used for the current numerical model. The flow behaviour was obtained from the isothermal compression tests conducted in the temperature range of 20−550∘C for EN AW-6082 and up to 1200∘C for 20 MnCr 5 at strain rates of 1, 10 and 50 s−1[21]. The flow curves were then implemented in the numerical
material model using the Hensel-Spittel approach to describe the material response within the relevant thermo-mechanical regime of the RFW process. Eq. 1 describes the flow stress kf as a function of plastic strain ε, temperature T and strain rate ε˙. The material specific model coefficients as defined by Wester et al. is shown in the Table 1 [21].

kf=Aem1*TTm8εm2e(m4ε+m6*ε)(1+ε)m5Tε˙(m3+m7T)(1)

Table 1. Material specific Hensel Spittel model coefficients [21]

Material | A | m₁ | m₂ | m₃ | m₄ | m₅ | m₆ | m₇ | m₈
20MnCr5 | 3.08e-8 | -0.0114 | 0.2787 | -0.0973 | -0.00053 | -0.00024 | -0.9644 | 0.00036 | 4.8115
EN AW-6082 | 1972.85 | -0.0111 | 0.0197 | -0.1779 | -0.00158 | 0.00243 | -0.7941 | 0.00087 | 0.1357

Methods - Experimental Investigation

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

The experimental study was carried out on a Genius Plus rotary friction welding machine from KUKA. Cylindrical specimens with a diameter of 40 mm and a length of 100 mm were used. The aluminium was positioned on the spindle (rotating) side, while the steel was mounted on the sledge (axially driven) side, Fig. 1(a), (b). Due to the significantly lower flow stress of aluminium, plastic deformation during friction welding occurred only in the aluminium component. Consequently, flash formation was confined to the aluminium specimen, while the steel remained essentially undeformed [22]. For this reason, temperature measurements were conducted exclusively in the steel specimen, where TC placement could be achieved without disturbing the material flow.

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Fig. 1. (a) Configuration and positioning of the steel and aluminium specimens during the welding process; (b) steel specimen with TCs prior to welding and (c) schematic illustration of TC positions with annotations.

Fig. 1. (a) Configuration and positioning of the steel and aluminium specimens during the welding process; (b) steel specimen with TCs prior to welding and (c) schematic illustration of TC positions with annotations.

Temperature evolution was recorded at 12 locations using K-type TC with a diameter of 2 mm . TC were arranged symmetrically, with even-numbered TCs located on one side of the specimen and odd-numbered TCs on the mirrored side providing six mirrored measurement positions to improve accuracy and repeatability. TC1 and TC2 were placed at the friction interface and inserted through holes oriented at 45∘ to the specimen axis in order to capture the interface temperature. The remaining 10 TCs were positioned along the axial direction at distances as shown in Fig. 1(c). These TCs were arranged symmetrically on both sides of the specimen. The specimen was machined according to the schematic shown in Fig. 1(c) to accommodate the TCs and to ensure consistent sensor positioning across all experiments.

Experiments were designed considering 5 welding parameters as independent variables including friction pressure, friction time, forging pressure, forging time and rotational speed, while the resulting temperature curves as the dependent variable. The experimental design, consisting of 15 PCs is shown in Table 2. The PCs were defined based on prior work by Dewidar et al., where the effect of parameters on the mechanical properties were investigated [22]. The current work extends these PCs by incorporating multiple rotational speeds and by limiting the friction time in accordance with machine constrains. This approach enables future comparison of the results and supports the extension of the process.

Table 2. The experimental design consisting of 15 PCs , with values in brackets representing force in kN

Parameter combination | Friction pressure in MPa | Friction time in s | Forging pressure in MPa | Forging time in s | Rotational speed in rpm
1 | 80 (100.5 kN) | 0.5 | 175 (220 kN) | 5 | 700
2 | 119.4 (150 kN) | 0.5 | 175 | 10 | 700
3 | 80 | 0.8 | 175 | 10 | 700
4 | 119.4 | 0.8 | 175 | 5 | 700
5 | 80 | 0.5 | 225 (282.7 kN) | 10 | 700
6 | 119.4 | 0.5 | 225 | 5 | 700
7 | 80 | 0.8 | 225 | 5 | 700
8 | 119.4 | 0.8 | 225 | 10 | 700
9 | 80 | 0.5 | 225 | 10 | 1,400
10 | 80 | 0.5 | 225 | 10 | 1,750
11 | 80 | 0.5 | 225 | 10 | 2,100
12 | 119.4 | 0.5 | 225 | 10 | 700
13 | 119.4 | 0.5 | 225 | 10 | 1,400
14 | 119.4 | 0.5 | 225 | 10 | 1,750
15 | 119.4 | 0.5 | 225 | 10 | 2,100

In addition to the in-situ temperature measurements, the axial force, displacement and rotational speed were recorded from the machine and can be used to define the mechanical and kinematic input parameters for the numerical model. Fig. 2(a) shows the recorded evolution of key process parameters for PC1 during the RFW cycle. For this experiment, the rotational speed was maintained at 700 rpm and a friction force of 100.5 kN was applied during the friction phase. At the end of this phase, the spindle rotation was stopped and the axial force was increased to 220 kN for the forging phase. Fig. 2(b) presents the resulting welded specimen with aluminium flash around the joint. The RFW tests were repeated two times for each PCs. Specimens from the PCs 5, 9, 10 and 11 were separated along the longitudinal axis by wire erosion to compare the flash geometries.

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Fig. 2. (a) Schematic evolution of the main process parameters of PC1, illustrating axial force-time and rotational speed-time relationships; (b) final friction-welded specimen after completion of the process.

Fig. 2. (a) Schematic evolution of the main process parameters of PC1, illustrating axial force-time and rotational speed-time relationships; (b) final friction-welded specimen after completion of the process.

Numerical Modelling

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

A 2D axisymmetric model was developed in the simulation software Simufact.forming v16 for the numerical analysis. The model consists of two cylinders, aligned coaxially along the same axis, as illustrated in Fig. 3. Material models were assigned to the workpieces according to the experimental arrangement. The body with the material model for aluminium was assigned as the workpiece and the steel as a deformable die. The upper, lower and support dies were defined as rigid bodies. A spring was used to substitute the axial force during the friction and the forging phase. The spring force was applied between the ground and the upper die making the upper die movable in the Zdirection.

[picture]

Fig. 3. Schematic representation of numerical model set up for 2D axisymmetric RFW simulation.

Fig. 3. Schematic representation of numerical model set up for 2D axisymmetric RFW simulation.

The combined friction model with a friction coefficient μ=0.05 and a friction factor m=0.05 was employed to describe the contact conditions between the two workpieces. These values were chosen based on preliminary simulations to match the experimental flow behaviour of the aluminium. This allows the model to capture the initial onset of material flow in aluminium without overestimating the interfacial resistance. These values can be further calibrated once the feasibility of the modelling approach is validated. In the simulation, particles were added to the same positions as the TCs on the steel workpiece to allow a direct comparison to the experiments. The model was set up in the cold forming/upsetting module in Simufact.forming. The geometries were discretised using elements of type Quads (10). To ensure that the numerical results are independent of the selected element size, a mesh dependency study was performed with global element sizes ranging from 3 mm to 0.5 mm . The computational time increased exponentially with the increasing mesh refinement as shown in Fig. 4(a). The average value of peak temperatures per increment during the friction phase in aluminium was evaluated for the different mesh sizes. The results showed that the predicted temperatures decreased with mesh refinement and stabilised at element sizes of 1 mm and 0.5 mm , which produced nearly identical results as in Fig. 4(b). A similar trend was observed for the average temperatures during the initial 2 seconds of the forging phase in the steel specimen, see Fig. 4(c). Considering the localized deformation at the joining interface, both the workpieces steel and aluminium were meshed with a global element size of 2 mm , with local mesh refinement applied at the weld interface to achieve an element size of 0.5 mm in this critical region, see Fig. 3. The required computational time and average temperatures when using the refinement box (RB) are also shown within the shaded regions in Fig. 4. Thus, a balance between computational efficiency and the accuracy in capturing the dominant thermo-mechanical trends of the process, consistent with the proof-of-concept nature of this study was acquired.

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Fig. 4. Results of mesh study representing (a) computation time as function of mesh size; (b) evolution of peak temperatures in aluminium during the friction phase; (c) average temperatures in steel during the initial 2 seconds of the forging phase.

Fig. 4. Results of mesh study representing (a) computation time as function of mesh size; (b) evolution of peak temperatures in aluminium during the friction phase; (c) average temperatures in steel during the initial 2 seconds of the forging phase.

The existing friction welding module within the Simufact.forming v16 has its limitations as stated in [8]. Therefore, a user-defined subroutine structure was implemented in Simufact.forming to account for contact-dependent thermal loading at the material interfaces. The application of heat flux using the subroutine substitutes the heat energy produced due to rotation enabling the 2D simulation and significantly reducing computational cost and time. The implementation relies on coordinated use of built-in user subroutine hooks executed at different stages of the solution procedure.

Fig. 5 represents a flowchart showing the subroutine calls in the Simufact. The subroutine UELOOP, which is called at every increment and iterates over all elements, was used to access the contact state information of the nodes of each element. For each increment, the contact identifiers (specifically contact touched body ID) associated with the nodes of the active elements were extracted and stored in nodal state variables. This approach allows the contact status of each node to be tracked throughout the simulation. In a subsequent step, the subroutine FORCDT, which is executed at every increment for a user-defined subset of nodes, was employed to apply thermal boundary conditions.

The subset of nodes evaluated by FORCDT was selected based on the corresponding body ID, thereby restricting the thermal calculations to the relevant material body. Nodes on which the frictional heat is to be applied were further identified and separated according to their specific contact identifiers.

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Fig. 5. Schematic representation of subroutine call in Simufact. Abbreviations: inc = increment, ielem = element number, numel = number of elements, inod = node number, nump = number of nodes.

Fig. 5. Schematic representation of subroutine call in Simufact. Abbreviations: inc = increment, ielem = element number, numel = number of elements, inod = node number, nump = number of nodes.

The frictional heat flux q, was then calculated following the approach described by Lei et al. [23] as stated in Eq. 2, where η is the mechanical efficiency factor (set to 0.9 ), τf is the interfacial friction shear stress and v is the friction velocity defined as the product of angular velocity ω and radius r.

q=ητfv(2)

The interfacial friction shear stress was further modelled using the combined friction model given by Eq. 3, where μ is the coefficient of friction, p is the normal pressure at the integration points of the interface elements, m is the friction factor and σf is the flow stress. The term μp represents the Coulomb friction model and the term mσf3 corresponds to the shear friction model. The μ and m values were provided as 0.05 to be consistent with the mechanical model. Taking the minimum of these two terms prevents the Coulomb friction stress from rising unrealistically with increasing normal pressure [23].

τf=min(μp,mσf3)(3)

The previously stored nodal state variables were accessed within the FORCDT to calculate and apply the heat flux to the interface nodes. In this way, the thermal loading was applied exclusively at the material interfaces where contact was detected. To verify the correct identification of contact nodes and the consistent application of the heat flux, the stored nodal state variables were visualized using the UPSTNO subroutine. This post-processing step enables direct monitoring of contact conditions and the correct execution of the user-defined logic throughout the simulation.

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

Fig. 6(a) presents exemplary the temperature histories recorded at all TC locations for PC10. This demonstrates the similarity between symmetrically positioned measurement points. It can be seen that the temperatures along the symmetrical counterpoints are in agreement confirming the assumption of rotational symmetry during the process. In RFW, the temperature is highest at the
contact surface, as measured by TC1 and TC2 and decreases progressively with increasing axial distance from the interface (TC3-TC12). The frictional heat causes a rapid plastic flow of the aluminium and in combination with the superimposed pressure, the aluminium can enter the channels of TC1 und TC2. This pushes the TCs back, meaning they no longer measure the temperature directly at the interface. As a result, temperature measurements obtained at these near-interface locations tend to exhibit higher variability and reduced reliability compared to the other TCs as shown in Fig. 6(b). The recorded near-interface measurements thus showed a higher failure rate during the experiments. The measuring points further away from the interface showed stable and reproducible measurement results, as the measuring position could be maintained throughout the entire test. Intermediate points were used to capture the axial temperature gradient, providing additional information for understanding the thermal field. Therefore, datasets TC4, TC10, TC6, TC12 and TC8 in the order of increasing distance from the weld interface respectively were used for the parameter study.

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Fig. 6. (a) Temperature histories measured at the twelve TC locations for PC10; (b) TC1 and TC2 showing deviating behaviour in PC15 caused by flash formation and local mechanical disturbances near the weld interface.

Fig. 6. (a) Temperature histories measured at the twelve TC locations for PC10; (b) TC1 and TC2 showing deviating behaviour in PC15 caused by flash formation and local mechanical disturbances near the weld interface.

The influence of RFW parameters on the peak temperature was evaluated using ANOVA and main effects analysis based on TC10. Although it is the second closest to the weld interface, TC10 provided reliable and reproducible measurements across all parameter combinations, making it the most representative of the thermal conditions governing joint formation. Fig. 7(a-e) presents the main effects plot for peak temperature at TC10, showing the adjusted mean temperature as a function of friction time, friction pressure, forging time, forging pressure and rotational speed. Friction time shows a strong positive effect on the peak temperature. An increase in friction time leads to a pronounced rise in the measured peak temperature. This behaviour is expected in RFW, as a longer friction phase increases the duration of interfacial sliding, thereby enhancing frictional heat generation and heat accumulation near the weld interface [ 10,11 ]. Friction pressure exhibits a relatively weak influence on the peak temperature at TC10. Increasing the friction pressure results in only a marginal increase in peak temperature. This indicates that, within the investigated range, friction pressure has a minor role in heat generation compared to other parameters. Forging time has minimal influence on the peak temperature. The main effects plot shows nearly identical mean temperatures for both levels of forging time, suggesting that peak temperature is largely established during the friction phase rather than the forging phase. Forging pressure demonstrates a moderate inverse effect on peak temperature. Increasing forging pressure slightly reduces the peak temperature. This reduction can be attributed to the enhanced plastic deformation and material expulsion during the forging phase, which promotes heat dissipation away from the interface. Rotational speed exhibits the strongest influence on peak temperature among all investigated parameters. A substantial increase in peak temperature is observed as rotational speed increases from the lowest to the highest level.

This trend reflects the direct relationship between rotational speed and frictional power input, as higher speeds significantly increase the rate of heat generation at the interface.

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Fig. 7. ANOVA main effects plots for peak temperature at TC10 as a function of the process parameters a) friction time, b) friction pressure, c) forging time, d) forging pressure and e) rotational speed.

Fig. 7. ANOVA main effects plots for peak temperature at TC10 as a function of the process parameters a) friction time, b) friction pressure, c) forging time, d) forging pressure and e) rotational speed.

To further isolate the influence of rotational speed, four PCs namely, 5, 9, 10 and 11 with identical process conditions and with increasing rotational speeds were grouped as Set 1 and compared. This enables a focused evaluation of the effect of rotational speed on peak temperature measured by the five TCs TC4, TC10, TC6, TC12 and TC8 positioned at progressively increasing distances from the weld interface respectively. Fig. 8(a-e) illustrates the variation of peak temperature with rotational speed for the TCs. For all locations, an overall increase in peak temperature with increasing rotational speed was observed. The results clearly demonstrate that the rotational speed influences peak temperature at all measured locations, but the effect diminishes with increasing distance from the weld interface (TC4 > TC10 > TC6 > TC12 > TC8). The strong temperature response at TC4 and TC10 further emphasizes the importance of rotational speed control for regulating the thermal conditions, which are critical for joint integrity in dissimilar material RFW. The temperature drop observed at 2,100rpm on TC4 falls within the experimental scatter due to the slipping out of TCs during the experiment.

The axial displacement of the sledge during the friction phase provides a quantitative measure of flash formation, as it directly reflects the plastic flow of aluminium driven by frictional softening at the interface. Measuring the sledge path in this stage allows the extent of material flow to be related to the frictional heat input, enabling assessment of how process parameters influence flash formation. Displacement data were exported from the machine and plotted as in Fig. 8(f) to examine trends with respect to rotational speed and friction pressure. Two sets of PCS were selected for the analysis. Set 1 consisted of the PC 5, 9, 10 and 11, all conducted at the lowest friction pressure of 80 MPa . Set 2 includes the PC 12, 13, 14 and 15, which employed the same sequence of rotational speeds but at the highest friction pressure of 119.4 MPa .

Set 1-PC5, PC9, PC10, PC11; | Set 2-PC12, PC13, PC14, PC15 |  |

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Fig. 8. (a-e) Peak temperature variation with rotational speed for TC4, TC10, TC6, TC12 and TC8; (f) axial displacement of the steel component during the initial friction phase for Set 1 (friction pressure 80 MPa ) and Set 2 (friction pressure 119.4 MPa ).

Fig. 8. (a-e) Peak temperature variation with rotational speed for TC4, TC10, TC6, TC12 and TC8; (f) axial displacement of the steel component during the initial friction phase for Set 1 (friction pressure 80 MPa ) and Set 2 (friction pressure 119.4 MPa ).

Fig. 8(f) presents the sledge path towards the aluminium counterpart during the initial friction phase. For both friction pressure levels, the sledge path increases with increasing rotational speed, indicating enhanced material softening and plastic flow due to higher frictional heat generation at elevated speeds. For example, in Set 1, increasing the rotational speed from 700 to 1,400rpm approximately doubles the sledge path, from 1.74 mm to 3.91 mm , while a further increase to 2,100 rpm results in a sledge path of 6.73 mm . At a certain rotational speed, the sledge path is consistently greater at higher friction pressure. With increasing friction pressure more aluminium is pushed out of the joining zone at lower temperatures and the associated higher flow stresses. Even when similar interface temperatures are reached at the same friction time, higher friction pressure promotes greater material extrusion, leading to an increased sledge path and flash formation. Fig. 9 illustrates the evolution of flash geometry with increasing rotational speed. A progressive increase in flash size is observed at higher rotational speeds, which correlates well with the increased sledge path recorded by the machine data during the friction phase. This agreement confirms the enhanced plastic flow at elevated rotational speeds. Further, the cross-sectional images show the penetrated aluminium into the drilled holes for TC1 and TC2, which caused the TCs to be pushed out of their original positions.

[picture]

Fig. 9. Wire eroded specimens of set 1 showing the effect of rotational speed on flash formation with (a) 700 rpm , (b) 1,400rpm, (c) 1,750rpm and (d) 2,100rpm (lowest friction pressure 80 MPa ).

Fig. 9. Wire eroded specimens of set 1 showing the effect of rotational speed on flash formation with (a) 700 rpm , (b) 1 , 400 r p m , (c) 1 , 750 r p m and (d) 2 , 100 r p m (lowest friction pressure 80 MPa ).

Rotational speed is the primary driver of interface temperature, as evidenced by the ANOVA results and the corresponding increase in sledge path with increasing speed. The enhanced thermal softening at higher rotational speeds leads to greater plastic flow during the friction phase. In contrast, the axial pressure exhibits a limited influence on peak temperature within the investigated range but plays a significant role in governing mechanical deformation. Excessive axial pressure can force softened material out of the joining zone resulting higher sledge path and flash formation. As temperature decreases away from the interface, material from the HAZ may be incorporated into the weld, potentially reducing bonding quality and compromising the integrity of the joining interface. Consequently, while rotational speed governs thermal input, axial pressure primarily modulates the extent of plastic upset and must be carefully balanced to optimise joint strength.

To evaluate the numerical model, Fig. 10 shows a comparison of the simulated results with the experimental measurement of PC11. In Fig. 10(a) the temperature histories at locations where post particles were added in the simulation corresponding to the positions of the embedded TCs in the steel specimen are compared. The numerical model qualitatively captures the overall temperature evolution at all measurement locations, including the rapid temperature rise during the friction phase, the occurrence of peak temperature and the subsequent cooling during the forging phase. The axial temperature distribution is also well reproduced, with the highest temperatures and steepest gradients at the weld interface, followed by a gradual decrease of the temperature with increasing distance from the interface.

Despite these qualitative agreements, significant deviations in the absolute temperatures were observed. At the weld interface (TC2), the predicted peak temperature is in good agreement with the experimental measurement, indicating that the overall magnitude of interfacial heat generation is reasonably captured by the numerical model. However, the simulated cooling rate is higher than that was observed in the experiment, suggesting an overestimation of heat dissipation from the interface region. The model underpredicts the peak temperatures, at other TCs particularly at TC4, TC10 and TC6, while predictions improve at locations farther from the interface at TC12 and TC8. The overall cooling trends were captured reasonably well. The underestimation of these temperature indicates that, although the current numerical model replicates the heat generation at the interface, it has not yet been fully calibrated. This is likely due to the assumptions in the interfacial friction coefficient and higher heat transfer coefficients which are difficult or impossible to measure. The experimental temperature data therefore provide a critical basis for the further refinement and calibration of the numerical model.

The comparison of flash geometries, as illustrated in Fig. 10(b), reveals that the experimentally obtained flash exhibits a stronger curvature and wraps back towards the aluminium component. In contrast, the simulated flash has a larger radius of curvature and does not fully close to the aluminium. Insufficient interface temperatures in the simulation results in higher flow stresses in the aluminium. This means that less material is displaced during the friction phase, resulting in reduced curvature in the flash. During the forging phase, the material flows more strongly in radial direction. This further supports the conclusion of uncertainties in the heat transfer coefficients including heat conduction to the aluminium component, heat partitioning at the interface and heat losses to the surroundings.

While the current model captures the main flow mechanisms and fundamental thermo-mechanical trends of RFW, quantitative agreement requires further refinement of the interfacial heat generation parameters and energy localization. In particular, using temperature data from the steel specimen to predict the flash geometry of aluminium remains challenging due to complex heat conduction and heat transfer at the dissimilar interface. The agreement between experiment and simulation confirms that the developed numerical model, implemented using a user-defined subroutine, provides a valid proof of concept for predicting temperature evolution and flash formation in dissimilar RFW.

[picture]

Fig. 10. (a) Temperature evolution for experimental PC 11 compared with numerical predictions; (b) corresponding comparison of flash geometries, illustrating differences in curvature and material flow between experiment and simulation.

Fig. 10. (a) Temperature evolution for experimental PC 11 compared with numerical predictions; (b) corresponding comparison of flash geometries, illustrating differences in curvature and material flow between experiment and simulation.

Summary and Outlook

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

An experimental study of RFW between 20 MnCr 5 and EN AW-6082 was performed to investigate the influence of process parameters on temperature evolution, plastic deformation and flash formation. The results indicate that rotational speed is the dominant factor controlling frictional heat generation, peak temperature and plastic flow, whereas friction pressure primarily affects plastic deformation with a limited influence on the peak temperature. The sledge path during the friction phase was found to be a reliable indicator of material softening and flash formation. In-situ temperature measurements in the steel specimen provided stable and reproducible thermal data, serving as a robust basis for numerical model calibration. Overall, this study demonstrates the key thermo-mechanical mechanisms governing steel-aluminium RFW and establishes a solid experimental foundation for future numerical modelling and process optimization of dissimilar metal joints.

The numerical simulations qualitatively reproduced the dominant trends in temperature, including peaks and axial distribution, as well as key features of flash morphology, such as curvature and material flow patterns. Quantitative deviations were observed, with the model predicting lower peak temperatures, higher cooling rates at the interface and thus predicting lower material flow. This highlights the importance of accurately modelling interfacial friction and heat transfer conditions. Overall, the combined experimental and numerical study demonstrates the fundamental thermomechanical mechanisms governing RFW and provides a platform for future model enhancement. The numerical model can be further developed to improve quantitative accuracy and readily extended to incorporate intermetallic formation behaviour, offering a predictive tool for a wide range of dissimilar metal joints.

Acknowledgements

The original version of this paper is available on https://www.scientific.net/MSF.1185.111.pdf

Funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) 252662854 - SFB 1153, subprojects A01 and B03.

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 127-138

doi: 10.4028/p-3pEJEy

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-16

Accepted:
2026-02-24

Online:
2026-04-15

Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

Mohamad El Mehtedi 1,a*, Pasquale Buonadonna 1, b, Barbara Reggiani 2,c, Riccardo Pelaccia 2, d, Lorenzo Donati 3,e, Gabriela Loi 1,f and Mauro Carta 1, g

1 Department of Mechanical, Chemical and Materials Engineering, University of Cagliari, 09123 Cagliari, Italy

2 DISMI Department of Sciences and Methods for Engineering, University of Modena and Reggio Emilia, Via Amendola 2, 42122 Reggio Emilia, Italy

3 DIN Department of Industrial Engineering, University of Bologna, Viale Risorgimento 2, 40136 Bologna, Italy

a* m.elmehtedi@unica.it, b  pasquale.buonadonna@unica.it, c  barbara.reggiani@unimore.it, d  riccardo.pelaccia@unimore.it, el.donati@unibo.it, f  gabriela.loi@unica.it, g  mauro.carta94@unica.it

*corresponding author

Keywords: roll bonding, aluminum alloy, magnesium sheets, formability, lightweight composites.

Abstract

With the increasing demand for lightweight materials, the combination of aluminum and magnesium sheets enables the development of advanced laminates with a balanced combination of strength and ductility, making them suitable for forming applications. This work investigates the effect of rolling temperature on the mechanical behavior and formability of AA1050/AZ31/AA1050 sheets produced by roll bonding in the temperature range of 250−450∘C. Tensile tests showed that the yield stress is weakly affected by rolling temperature, whereas the ultimate tensile strength increases up to 350∘C and then stabilizes. The elongation at fracture increases monotonically with temperature, indicating improved ductility at higher rolling temperatures. Microhardness measurements revealed softening of the aluminum sheets with increasing temperature, while limited variations were observed in the AZ31 sheet. Formability was evaluated by Erichsen Cupping test. The maximum load and extension at break remained nearly constant over the investigated temperature range; however, higher rolling temperatures led to reduced delamination and improved interfacial bonding integrity during deformation. The results indicate that roll bonding at elevated temperatures promotes better strain distribution and enhanced bonding quality. Overall, roll bonding at 450∘C provides the most favorable combination of mechanical performance, formability, and interfacial stability, making the produced sheets suitable for lightweight forming applications.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

The increasing demand for lightweight structural materials with enhanced mechanical performance has driven significant research efforts toward the development of metal-metal sheets composites; aluminum and magnesium are two key lightweight metals widely employed to achieve weight reduction [1].
In general, aluminum alloys combine low density with good corrosion resistance, high thermal and electrical conductivity, and excellent formability, making them highly versatile materials for structural and semi-structural applications. Their mechanical properties can be tailored over a wide range through alloying and thermomechanical treatments (including heat treatments and work hardening) [2,3], allowing a balance between strength, ductility, and manufacturability.
Magnesium alloys, on the other hand, are some of the lightest structural metallic materials currently in use and are valued for their high specific strength and stiffness [4]. These characteristics make them particularly attractive for aggressive lightweight design. However, their broader adoption is limited by intrinsic drawbacks, such as reduced ductility and formability at room temperature, pronounced anisotropy, and generally lower corrosion resistance compared to aluminum alloys due to its higher chemical reactivity [5,6]. As a result, significant research efforts are focused on
overcoming these limitations through alloy design, processing innovations, and the development of hybrid or composite material solutions [7].
Roll Bonding (RB) is an effective solid-state joining process for the fabrication of layered metal composites [8], in which two or more metal sheets are stacked and plastically deformed together by rolling, promoting bonding through severe plastic deformation at the interface. Prior to rolling, the sheet surfaces are typically cleaned and mechanically roughened to remove oxide layers and enhance intimate contact [9,10]. When repeated through multiple cycles with a 50% reduction, Roll Bonding can be extended to Accumulative Roll Bonding (ARB), a Severe Plastic Deformation (SPD) technique capable of producing multilayer composites with refined or ultrafine-grained microstructures and enhanced mechanical properties, first introduced by Saito et al. [11].
Early studies have shown that severe plastic deformation induced by ARB can produce high-quality Al/Mg interfaces and refined microstructures, leading to improved mechanical properties with increasing ARB cycles. These improvements can generally be attributed to work hardening and grain refinement in both constituents [12,13]. Temperature and reduction percentage strongly affect interlayer bonding; an increase in either parameter, leads to improved bonding [14].
However, the evolution of the mechanical properties in Al/Mg composites is non-linear; for instance, it has been observed that while strength initially increases due to strain hardening and grain refinement, it may decrease in later cycles (typically after the 3rd or 4th pass) due to the loss of structural integrity in the magnesium layers and the presence of some intermetallic compounds (IMCs) [15].
More recent investigations have explored modified ARB routes, such as asymmetric accumulative roll bonding (A-ARB), which introduces additional shear strain through different roll speeds. This approach has been shown to improve strain distribution and further refine the grain structure, potentially enhancing the mechanical bonding at the interface [16]. Furthermore, post-processing heat treatments (annealing) have been studied to mitigate the high internal stresses induced by ARB. While annealing at moderate temperatures (around 200∘C ) can improve ductility through recovery and recrystallization of the Al and Mg layers, higher temperatures rapidly accelerate IMCs growth, which can negatively affect the composite's overall performance [17].
Despite the results shown in literature, most existing studies primarily address microstructural evolution and uniaxial tensile behavior. However, in industrial applications, sheet materials are rarely utilized in their flat, as-rolled state; they are almost always subjected to secondary forming processes involving complex stress states. Consequently, formability assessments such as the Erichsen cupping test are of paramount importance to evaluate the material's suitability for real-world manufacturing. Specifically, the role of rolling temperature in controlling the onset of interfacial damage and the eventual fracture of the protective cladding during forming remains a critical gap in current research. The present work aims to bridge this gap by investigating AA1050/AZ31/AA1050 composite sheets produced by roll bonding. By systematically correlating tensile behavior, fracture morphology, and interfacial stability, this study provides new insights into the processing-deformation-formability relationship using Erichsen cupping test.

Materials and Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

The materials investigated in this study are a commercially pure aluminum alloy AA1050 ( 99.5wt% Al min .) and the AZ31 magnesium alloy, which contains approximately 3wt%Al and 1wt%Zn. AA1050 was selected for its excellent corrosion resistance, high ductility, and good workability, while AZ31 was chosen due to its low density and high specific mechanical strength. The combination of these two alloys enables the production of lightweight multilayer sheets with complementary physical and mechanical properties.
The AA1050 alloy was supplied in sheet form with a thickness of 1.2 mm , whereas the AZ31 alloy was supplied in sheet form with a thickness of 0.8 mm . Both materials were cut to identical dimensions of 100 mm in length and 50 mm in width prior to processing. The thickness ratio ( 1.2 mm Al/0.8 mmMg ) was selected based on industrial requirements to ensure process stability during roll bonding, improved bonding quality, and protection of the Mg core from oxidation and cracking.

While a higher Mg fraction could increase specific strength, it would reduce formability and increase the risk of interfacial damage due to the lower ductility of AZ31. The chosen configuration therefore represents an industrially driven compromise between manufacturability, interfacial integrity, and mechanical performance. Future work will specifically address the optimization of the Al/Mg thickness ratio to further improve lightweight efficiency and strain partitioning.

Roll Bonding Procedure

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

Multilayer composite sheets were produced by Roll Bonding (RB). Prior to bonding, the surfaces of the AA1050 and AZ31 sheets were mechanically treated by paper grit 80 to remove surface oxides and contaminants and to increase surface roughness, thus enhancing interfacial bonding. The AZ31 magnesium sheets were ground on both surfaces, while the AA1050 aluminum sheets were ground only on the surface intended to be in contact with the magnesium sheet.
The sheets were stacked in an AA1050/AZ31/AA1050 configuration and secured with steel wire at both ends of the stack. Then, the assembled sheet stack was subjected to a preheating treatment in a furnace. Preheating temperatures ranged from 250∘C to 450∘C, with increments of 50∘C, and a holding time of 10 minutes was adopted for each temperature condition. Immediately after preheating, the sheet stack was extracted from the furnace and rapidly transferred (about 2 sec ) to the rolling mill to minimize heat loss due to environmental exposure (Figure 1).

[picture]

Fig. 1. Scheme of the Roll Bonding process for AA1050/AZ31 composite sheets.

Fig. 1. Scheme of the Roll Bonding process for AA1050/AZ31 composite sheets.

Considering an initial total thickness of 3.2 mm , the rolling gap was set to 1.6 mm , corresponding to the theoretical final thickness after rolling (50% reduction).
Rolling was performed without the use of lubricants. After rolling, the composite sheets were aircooled to room temperature.

Characterization Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

Specimens for microstructural and mechanical characterization were extracted from the rolled composite sheets. Cross-sectional samples were mounted using cold-setting resin, followed by standard grinding and polishing procedures to obtain surfaces suitable for optical microscopy in order to study the thickness distribution of each layer at different rolling temperatures. Optical microscopy was employed to investigate the multilayer architecture of the composite sheets, focusing on layer thickness, interface continuity, and macroscopic bonding quality. The thickness of each layer was measured at 15 locations per layer to ensure statistical reliability (Figure 2.a).
Tensile specimens were machined from the composite sheets with the loading direction parallel to the rolling direction (RD). Figures 2.b-c show the tensile samples obtained from the composite sheets and the geometry of the tensile sample respectively.

[picture]

Fig. 2. (a) Cross-sectional image of a representative sample showing multiple thickness measurements of each layer, with the rolling direction aligned with the scale bar; (b) photograph of the specimens after machining, prior to mechanical testing; and (c) geometry and main dimensions of the dog-bone specimen.

Fig. 2. (a) Cross-sectional image of a representative sample showing multiple thickness measurements of each layer, with the rolling direction aligned with the scale bar; (b) photograph of the specimens after machining, prior to mechanical testing; and (c) geometry and main dimensions of the dog-bone specimen.

Mechanical characterization was carried out by tensile testing to determine the yield strength ( σ0.2 ), ultimate tensile strength (UTS), Young's modulus (E), and elongation at fracture (A%). Tests were performed using an Instron 5585 H universal testing machine equipped with an Instron 2620-601 extensometer, applying a constant displacement rate of 1 mm/min.
Formability tests were also conducted to evaluate the deformation behavior of the composite sheets under bending conditions. Erichsen Cupping tests were performed using the same universal testing machine employed for tensile testing and were carried out with the aid of the mechanical setup with the tools shown in Figure 3.a-c. The system consisted of two clamping rings, a tightening device, and a hemispherical punch ( 19 mm diameter), all supported by two plates required for mounting the setup on the testing machine. For each rolling condition, three circular specimens with a diameter of approximately 50 mm were prepared. The tests were carried out by placing a Teflon sheet between the punch and the specimen to reduce friction and applying a constant punch displacement rate of 2 mm/min. Load and displacement data were continuously recorded during all tests.
After testing, the specimens were cut along one diameter to examine the bonding between all layers and subsequently observed by optical microscopy.

[picture]

Fig. 3. Experimental setup for formability testing: (a) mechanical tooling installed on the universal testing machine, (b) three-dimensional view of the tooling assembly, and (c) dimensioned schematic cross-section of the system.

Fig. 3. Experimental setup for formability testing: (a) mechanical tooling installed on the universal testing machine, (b) three-dimensional view of the tooling assembly, and (c) dimensioned schematic cross-section of the system.

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

Figure 4 shows the effect of rolling temperature on the thickness distribution of the individual AA1050 and AZ31 sheets, as well as on the total composite sheet thickness. The total thickness remains close to the nominal value of 1.6 mm for all processing conditions, confirming good control of the rolling process. However, the thick partitioning between aluminum and magnesium sheets clearly depends on rolling temperature. At lower rolling temperatures, the AA1050 sheets accommodate a larger fraction of the imposed deformation, while the AZ31 sheet remains relatively thicker. With increasing rolling temperature, the thickness of the AA1050 sheets remains nearly constant up to 400∘C and then decreases at 450∘C, corresponding to a higher deformation level with respect to the initial thickness. Conversely, the thickness of the AZ31 sheet remains approximately constant up to 400∘C, then increases at 450∘C. This indicates a reduced contribution to plastic deformation at the highest temperature. This opposite trend reflects the different temperature dependence of the flow stress of the two materials [18,19]. At different temperatures, the flow stresses of AA1050 and AZ31 differ, leading to a redistribution of strain between the sheets and a modification of deformation partitioning during roll bonding.

[picture]

Fig. 4. Effect of rolling temperature on sheet thickness after Roll Bonding: (a) thickness evolution of individual AA1050 and AZ31 sheets; (b) total composite sheet thickness compared with the nominal thickness.

Fig. 4. Effect of rolling temperature on sheet thickness after Roll Bonding: (a) thickness evolution of individual AA1050 and AZ31 sheets; (b) total composite sheet thickness compared with the nominal thickness.

Table 1. Mechanical properties of composite sheets.

Rolling temperature | UTS [MPa] | σ0.2 [MPa] | A% [%] | Young Modulus [GPa]
250°C | 115 ± 21 | 114 ± 21 | 0.5 ± 0.1 | 50 ± 1
300°C | 141 ± 5 | 121 ± 7 | 0.7 ± 0.0 | 64 ± 2
350°C | 153 ± 5 | 116 ± 15 | 6.1 ± 0.1 | 80 ± 2
400°C | 154 ± 1 | 118 ± 5 | 9 ± 0.5 | 62 ± 12
450°C | 153 ± 2 | 117 ± 3 | 15.3 ± 0.7 | 58 ± 2
AA1050 - base material | 70 ± 1 | 40 ± 2 | 36.4 ± 1.2 | 69 ± 1
AZ31 - base material | 223 ± 2 | 130 ± 1 | 7.6 ± 2.0 | 43 ± 2

[picture]

Fig. 5. (a) Representative engineering stress-strain curves of AA1050/AZ31/AA1050 specimens processed at different rolling temperatures together with the base materials (AA1050 and AZ31); (b)lateral view of fracture surfaces of the tensile specimens tested at different rolling temperatures.

Fig. 5. (a) Representative engineering stress-strain curves of AA1050/AZ31/AA1050 specimens processed at different rolling temperatures together with the base materials (AA1050 and AZ31); (b)lateral view of fracture surfaces of the tensile specimens tested at different rolling temperatures.

Figure 5.a shows the engineering stress-strain curves of the AA1050/AZ31/AA1050 composite sheets produced at different rolling temperatures while Table 1 and Figure 6 summarize the corresponding mechanical properties.
Rolling temperature has a limited influence on yield stress, which remain within the same range for all investigated conditions. In contrast, ultimate tensile strength (UTS) and elongation at fracture are strongly affected by rolling temperature. The UTS increases with temperature up to approximately 350∘C, reaching values close to 153 MPa , and then remains nearly constant at higher temperatures. At the same time, the elongation at fracture increases markedly with rolling temperature, rising from values below 1% at 250−300∘C to more than 15% at 450∘C. This combination of stable strength and strongly enhanced ductility indicates a significant improvement in deformation compatibility at elevated rolling temperatures. While the tensile specimens used in this work have relatively small gauge dimensions (Figure 2), which can slightly overestimate the measured ductility, Nie et al. [13] reported similar mechanical properties for Al/Mg/Al composites after the first ARB cycle at 400∘C. Temperatures higher than 450∘C are not expected to further improve the mechanical response of the laminate. In this temperature range, the AZ31 layer experiences extensive dynamic recrystallization accompanied by significant grain growth and softening, which progressively lowers the flow stress and promotes slip-dominated deformation, thereby reducing its effective load-bearing capability and leading to unfavorable strain partitioning between the layers [20,21].
The relatively large deviations observed in the Young's modulus values can be attributed to the intrinsic characteristics of multilayer composites and the experimental methodology. First, the AA1050/AZ31/AA1050 laminate is mechanically heterogeneous, and small variations in layer thickness, local bonding quality, and strain partitioning can significantly affect the initial elastic slope measured during tensile testing. Furthermore, the elastic response of the composite is governed by the rule of mixtures between aluminum and magnesium layers, whose elastic moduli differ substantially, amplifying the scatter when minor geometric or interfacial variations are present.
The fracture surfaces shown in Figure 5.b provide direct evidence of this transition. At the lowest rolling temperature ( 250∘C ), clear delamination between the AA1050 and AZ31 layers is observed, accompanied by limited necking and a relatively flat fracture profile. This indicates insufficient interfacial bonding and premature interfacial failure under tensile loading. As the rolling temperature increases, the extent of interfacial damage is progressively reduced. At 450∘C, no evident delamination is observed, and the fracture surface is characterized by pronounced necking and a more ductile morphology even in AZ31 layer, confirming the formation of a stronger metallurgical bond between the layers.
In addition, a brief consideration of the interfacial shear strength is necessary to better interpret the mechanical response of the AA1050/AZ31/AA1050 laminates. In roll-bonded dissimilar composites, the overall load transfer between layers is strongly controlled by the interfacial shear strength, which depends on oxide layer disruption, plastic deformation at the interface, and metallurgical bonding quality. The results obtained in this study suggest that increasing the rolling temperature enhances the interfacial shear strength, as evidenced by the progressive reduction of delamination and the more cooperative deformation of the layers during tensile loading. From a processing perspective, further improvements in shear strength could be achieved through higher reduction ratio, the adoption of asymmetric rolling to introduce additional shear strain, or post-rolling heat treatments aimed at promoting diffusion bonding while limiting excessive intermetallic growth [22,14].

[picture]

Fig. 6.Variation of the mechanical properties of sheets composites with processing temperature: (a) Young's modulus, (b) yield stress, (c) maximum strain at fracture, and (d) maximum stress.

Fig. 6. Variation of the mechanical properties of sheets composites with processing temperature: (a) Young's modulus, (b) yield stress, (c) maximum strain at fracture, and (d) maximum stress.

The formability of the composite sheets was evaluated by Erichsen cupping tests, and representative results are shown in Figure 7 and Figure 8. Figure 7.a illustrates the typical load-extension curves obtained for specimens processed at 250∘C,350∘C and 450∘C, along with the final shapes of the deformed sheets after testing. It shows that the curves for different rolling temperatures exhibit a similar overall trend, characterized by a progressive increase in load up to a maximum value, followed by fracture.
Figure 7.b provides a clear correlation between the global mechanical response measured during the Erichsen cupping test and the damage mechanisms observed in the deformed composite sheets. The maximum load and extension values are comparable, confirming that the global formability response is only weakly affected by rolling temperature. However, the cross-sectional views reported in Figure 7.b reveal significant differences in deformation and damage evolution. At the lowest rolling temperature ( 250∘C ), the specimen exhibits pronounced interfacial separation and localized cracking, indicating insufficient bonding strength to sustain large plastic deformation. At the intermediate temperature ( 350∘C ), interfacial integrity is improved, although local damage is still visible near the apex. In contrast, the specimen processed at 450∘C shows a more uniform deformation profile with limited delamination, demonstrating enhanced metallurgical bonding and improved strain compatibility between the aluminum and magnesium sheets. These observations highlight that similar load-extension responses may mask substantial differences in damage mechanisms. While global formability parameters remain nearly constant, the rolling temperature strongly influences interfacial integrity and failure mode. The improved performance observed at higher temperatures can therefore be attributed to enhanced bonding quality, which delays interfacial failure and allows more homogeneous deformation during the Erichsen cupping test.

[picture]

Fig. 7. Results of the formability test at 450∘C : (a) load-extension curve recorded during the test, and (b) cross-section view of the deformed specimen after forming.

Fig. 7. Results of the formability test at 450 ∘ C : (a) load-extension curve recorded during the test, and (b) cross-section view of the deformed specimen after forming.

As reported in Figure 8 and Table 2, both the maximum load and the extension at break remain relatively stable over the entire investigated temperature range. The extension at break varies between approximately 4.1 and 4.8 mm , while the maximum load is around 3000 N . This indicates that the global response measured during the Erichsen cupping test is only weakly dependent on rolling temperature.
Despite the similar global response, macroscopic examination of the tested specimens reveals a clear influence of rolling temperature on damage evolution. Sheets produced at lower rolling temperatures exhibited early interfacial separation and localized cracking, whereas specimens rolled at higher temperatures showed reduced delamination and more uniform deformation. This observation indicates that improved metallurgical bonding at elevated rolling temperatures enhances formability by delaying interfacial failure and delamination.

[picture]

Fig. 8. Formability test results as a function of temperature: (a) extension at break and (b) maximum load.

Fig. 8. Formability test results as a function of temperature: (a) extension at break and (b) maximum load.

The correlation between rolling temperature and interfacial integrity observed in our Erichsen cupping tests is further supported by the findings of Abbasi and Sajjadi [23], who investigated the three-point bending behavior of Al/AZ31/Al composites. In their study, low rolling temperatures (200∘C) resulted in multiple "load drops" during bending, which were attributed to sequential delamination and the inability of the magnesium core to sustain tensile stresses at the outer radius. Similarly, our specimens processed at 250∘C exhibited early interfacial separation and localized cracking during cupping (Fig. 6b), confirming that a low-temperature mechanical bond is insufficient
for complex stress states. In summary, sheets processed at lower temperatures exhibit early interfacial separation during forming, whereas specimens rolled at higher temperatures show more uniform deformation and stable interfaces This demonstrates that effective formability is primarily governed by interfacial integrity rather than by global stiffness or strength.

Table 2. Formability test results of AA1050/AZ31 composite sheets produced at different preheating temperatures.

Temperature [°C] | Maximum Load [N] | Extension to Break [mm]
250 | 2679 | 4.1
300 | 3339 | 4.8
350 | 2975 | 4.4
400 | 2910 | 4.6
450 | 2808 | 4.5

Conclusion

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

AA1050/AZ31/AA1050 composite sheets were successfully produced by roll bonding in the temperature range of 250−450∘C, demonstrating the feasibility of combining aluminum and magnesium into lightweight sandwich structures for forming applications. The main conclusions are as follows:

rolling temperature was identified as a key processing parameter controlling interfacial quality and overall mechanical behavior.

The ultimate tensile strength (UTS) increased with rolling temperature up to ~350∘C before stabilizing, while elongation at fracture showed a continuous improvement.

Erichsen cupping tests revealed that global formability parameters (maximum load and extension at break) were only weakly affected by rolling temperature; however, higher temperatures significantly reduced interfacial delamination and promoted more uniform deformation across the layers.

Roll bonding at 450∘C provided the best compromise between strength, ductility, formability, and interfacial integrity, making the resulting sheets suitable for lightweight forming applications.
Future work will focus on optimizing the initial thickness ratio between AA1050 and AZ31 sheets to tailor strain partitioning and improve the mechanical response. Further investigations will include determining forming limit diagrams (FLDs) through additional tests under different strain paths to better define the forming window of the composite. Process modeling will also be developed to support the optimization of roll-bonding parameters for targeted lightweight applications.

Funding

The original version of this paper is available on https://www.scientific.net/MSF.1185.127.pdf

This study is funded by the Italian Ministry of University and Research, PNRR, Mission 4, Component 2, Investment line 1.1-Call for tender No. 104 published on 2.2.2022 by the Italian Ministry of University and Research (MUR), funded by the European Union-NextGenerationEUProject Title: "SCULPTROL-SCULPTuring by ROLling for multi-material bonding"-CUP: F53D23001830006-Grant Assignment Decree No 961 adopted on 30 June 2023 by the Italian Ministry of Ministry of University and Research (MUR).

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 139-147

doi: 10.4028/p-D4Bxr8

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-12

Accepted:
2026-02-24

Online:
2026-04-15

Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Lasse Malaske 1,2,a*, Uceu Suhuddin 2, b and Benjamin Klusemann 1,2,c

¹Institute for Production Technology and Systems, Leuphana University Lüneburg, Universitätsallee 1, 21335, Lüneburg, Germany

2 Solid State Materials Processing, Institute of Material and Process Design, Helmholtz-Zentrum Hereon, Max-Planck-Straße 1, 21502, Geesthacht, Germany

a* lasse.malaske@leuphana.de, b uceu.suhuddin@hereon.de, c benjamin.klusemann@leuphana.de

*Corresponding Author: lasse.malaske@leuphana.de

Keywords: aluminum, titanium, solid-state, magnesium, silicon, refill friction stir spot welding, dissimilar, welding, interface, surface preparation, mechanical properties.

Abstract

Optimizing the performance and reliability of welding techniques for dissimilar aluminum (Al) to titanium (Ti) is a promising way to establish new applications in aerospace industry. Due to structural weight reduction, lightweight materials can help to minimize fuel consumption and save emissions. Solid-state welding technologies allow short joining cycles and metallurgical changes, residual stresses and severe intermetallic compound formation can be reduced by limited thermal exposure. Besides temperature and plastic deformation, intimate contact plays an important role for diffusion. In this work, AlMgSi alloys with systematic variations of Mg and Si alloying elements, were welded to Ti6Al4V (Ti64) by refill Friction Stir Spot Welding. The focus lays on the effect of Ti64 sheet surface roughness, varied by different surface preparations. Additionally, the influence of the plunge depth, the distance between the tool and the Ti64 sheet surface is analyzed. It was found that a reduced tool to interface spacing has a beneficial influence on joint integrity. Grinding trenches allowed better bonding compared to the pit-like surface structure generated by sandblasting, which led to an increase in mechanical lap-shear properties. Knurling the grinded surfaces resulted in high standard deviation, as most likely not the whole interface area was bonded. However, the partially outstanding properties showed that a beneficial effect can be expected due to mechanical interlocking mechanisms, when sufficient diffusion is ensured.

Introduction

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Lightweight materials such as aluminum (Al) and Titanium (Ti) are widely used in the aerospace industry. While Ti is characterized by outstanding high temperature resistance and strength, Al is cost effective. To balance the structural and economic demands, Al/Ti hybrid structures are of great interest [1]. However, techniques for joining them efficiently are still under investigation, as conventional fusion welding techniques are associated with issues based on the huge material property discrepancies of Al and Ti [2].
Solid-state welding, such as friction stir welding and refill Friction Stir Spot Welding (refill FSSW), represents excellent solutions and has been proven to be suitable for several alloy combinations [3]. A crucial factor for obtaining sound joints is the plunge depth (PD). It was found for refill FSSW of AA5754 to galvanized DP600 that a higher PD up to close to the interface allowed for a strong increase in mechanical joint properties [4]. In terms of welding dissimilar Al/Ti, challenges occur as the formation of brittle Al3Ti [5] and the introduction of residual stresses due to the substantial differences in physical and chemical properties of Al and Ti[1]. These difficulties can lead to interfacial fractures, often occurring well below the strength of the Al , for instance in a friction stir welded ADC12 Al/Ti joint [6], failing at 62% of the base material.
To improve joint strength, surface pre-treatments may be a suitable approach to maximize the effect of mechanical interlocking mechanisms at the interface to support the diffusion-based bonding. Additionally, it was shown during the solid-state deposition process Friction Surfacing, that
increasing topography depth led to a decreased average grain size [7]. This could potentially support diffusion mechanisms due to higher grain boundary density [8] and therefore accelerate the bond formation, which could allow to reduce the process duration and correlated thermal exposure. Enhancements in the mechanical properties were shown for other dissimilar material combinations, such as AA5052 resistance spot welded to galvanized steel [9] and AA5182 laser welded to PA66 [10]. Particularly the increased surface roughness and correlated with that the increased interfacial area for diffusion, has shown positive effects on the weld strength [9,11]. For ultrasonic spot welding of AA1100 to stainless steel [12], it was found that cleaned stainless steel surfaces exhibited better mechanical properties compared to electrolytically polished surfaces, indicating the relevance of surface roughness. Nonetheless, surface roughness can also have a negative influence, as e.g. within grinding trenches small voids and unbounded regions can occur which then act as crack initiation sites [13,14]. In this context, He et al. [15] showed that ordered patterns enable superior improvements compared to simple grooves. Perpendicular and 45∘ diagonal textures provided better strength under multiaxial loading due to improved mechanical interlocking and microstructural pinning effects [15]. In this study, the effect of surface roughness and topography of Ti64 as well as PD is investigated during dissimilar Al/Ti welding using refill FSSW, to vary micro-mechanical interlocking and diffusion bonding mechanisms. Two different PD levels and four different Ti64 sheet surface pretreatments were tested to allow for a better understanding of the influence on joint integrity including mechanical properties and failure mode.

Materials and Methods

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Joints of Ti64 ( 100×25.4×2.0 mm3 sheets) with five different Al alloys ( 100×25.4×1.8 mm3 sheets), based on AA6013 with systematic varied Mg and Si content, were welded via refill FSSW. The Al alloys were gravity die casted at 700∘C, water quenched and afterwards the billets were extruded at 450∘C. Optical emission spectroscopy was used to determine the chemical composition of the alloys, see Table 1.

Table 1. Chemical composition of the five Al alloys determined by optical emission spectroscopy, based on AA6013 ( Al1Mg1Si ). For the other alloys, the amount of Mg and Si was systematically varied. All values in wt.%.

Alloy | Si | Mg | Cu | Fe | Mn | Zn | Cr | Ti | Al
Al1Mg1Si | 0.98 | 1.10 | 0.6 – 1.1 | Max 0.5 | 0.2 – 0.8 | Max. 0.25 | Max. 0.1 | Max. 0.1 | Res.
Al1Mg0Si | 0.01 | 1.06 | Res.
Al1Mg1.5Si | 1.52 | 1.03 | Res.
Al3Mg1Si | 1.10 | 3.13 | Res.
Al5Mg1Si | 0.78 | 5.30 | Res.

Refill FSSW, developed by Helmholtz-Zentrum Hereon [16], is a solid-state spot welding technique that allows short welding cycles and the production of energy efficient joints. The process is schematically presented in Figure 1 and works with a tool consisting of three components. The outer clamping ring (CR) is stationary and applies a constant force to the sheets below to keep these in position. The inner components are the probe in the center and the shoulder in between the probe and the CR. Both rotate, typically at the same rotational speed and direction, whereby the shoulder plunges into the workpiece and the probe retracts in the opposite direction. Due to the friction between the tool and workpiece, the upper Al sheet heats up and plasticizes. The plasticized Al deforms and flows into the cavity which is created by the probe during the first process stage. Then an optional dwell time can be used to extend the process without any translational, but still rotational tool movement. Afterwards shoulder and probe retract to the original positions and the process finished with lifting all 3 components from the workpiece. Each joint was welded in overlap configuration with the Al sheet being placed above the Ti64 sheet, with an overlapping area of 25.4×25.4 mm2. The PD of the shoulder into the Al sheet varied between 1.2 and 1.5 mm , meaning between 0.6 and 0.3 mm distance of the shoulder to the Ti64 surface. Except for the PD, all other process parameters
were kept constant in this study and were selected as: dwell time (DT) of 5 s , rotational speed (RS) of 2200 rpm , clamping force (CF) of 8 kN and plunge time of 1 s .

[picture]

Fig. 1. Schematic illustration of the refill FSSW process, showing the tool components, sheet orientation and process stages:1) Plunging phase, 2) Dwelling phase, 3) Retraction phase, 4) Tool removal.

Fig. 1. Schematic illustration of the refill FSSW process, showing the tool components, sheet orientation and process stages:1) Plunging phase, 2) Dwelling phase, 3) Retraction phase, 4) Tool removal.

Lap shear testing was performed at room temperature using a Zwick-Roell 1478 universal testing machine. A constant crosshead speed of 1 mm/min was used, and the sample dimensions are presented in Figure 2. The topography micrographs of the surface pre-treatments and after lap-shear testing also the fracture surface micrographs were investigated by using an FEI Quanta 650 scanning election microscope (SEM). Additionally, the pre-treatment surface roughness was analyzed by a profilometer (Keyence VR 5200) using the average of 10-line measurements.

[picture]

Fig. 2. Schematic of the lap-shear test specimens for determination of the mechanical properties. All dimensions in mm.

Fig. 2. Schematic of the lap-shear test specimens for determination of the mechanical properties. All dimensions in mm.

The surface pre-treatment of the Al sheet on both surfaces was kept constant using #320SiC paper, while the surface of the Ti64 sheet was pre-treated in different ways, see Table 2. Grinding the sheets was performed immediately prior to welding while sandblasting (SB) as well as knurling (KN) were performed in advance, handling the Ti64 sheets in ethanol afterwards to suppress surface oxidation.

Table 2. Description of the surface preparation methods used on the Ti64 surface.

Method | Designation | Tool / Medium
Grinding (#100) | #100 | SiC paper #100
Grinding (#320) | #320 | SiC paper #320
Sandblasting | SB | Al₂O₃, 100-150 μm
Grinding (#320) and Knurling | KN | SiC paper #100; Knurling tool 290196, 20 × 8 / 0.5

SB was used to remove the oxide layer of the Ti64 and to generate a homogenous surface structure that varies clearly from the surface structure after grinding to change the contact conditions of the Al and Ti64 sheets in the interface, Figure 3a,b. As blasting medium Corundum ( Al2O3, mesh size 100150μ m ) was selected and the parameters were 60 s of SB, 3 bar of pressure at 300 mm distance in a 30∘ angle.
For generating indents that may allow macro-mechanical interlocking, knurling was performed using a knurling tool (290196, 20×8/0.5 ) with a tooth pitch of 0.5 and knurl type GE 30∘ (Hoffmann GmbH Qualitaetswerkzeuge, Munich, Germany), Figure 3c. A traverse speed of 2 mm/s and a tilt angle of 0∘ was used. Prior to knurling, the sheets were carefully ground using #320 SiC paper to remove the oxide layer.

Table 3. Overview of the generated surface topographies, depending on the individual surface pre-treatment.

Method | Rz (Avg.) | Rz (Std. Dev.) | Ra (Avg.) | Ra (Std. Dev.)
Grinding #100 | 5.99 | 0.852 | 0.844 | 0.105
Grinding #320 | 4.80 | 0.882 | 0.689 | 0.156
Sandblasting | 6.38 | 2.64 | 0.803 | 0.116
Knurling | 92.27 | 18.458 | 18.647 | 4.766

[picture]

Fig. 3. Ti64 surface appearance after pre-treatment a) SEM micrograph of grinded Ti64 using SiC paper #100, b) SEM micrograph of SB Ti64 surface, c) Topography measurement of the KN pattern, d) sketched orientation of the surface pre-treatments with regard to the lap-shear testing direction.

Fig. 3. Ti64 surface appearance after pre-treatment a) SEM micrograph of grinded Ti64 using SiC paper #100, b) SEM micrograph of SB Ti64 surface, c) Topography measurement of the KN pattern, d) sketched orientation of the surface pre-treatments with regard to the lap-shear testing direction.

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Tool to interface spacing.

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

All joint combinations and surface pre-treatments performed at 1.5 mm PD resulted in high-quality joints with sound mechanical properties. Figure 4 presents the results, revealing that PD reduction of 0.3 mm , i.e. PD of 1.2 mm , resulted already in a significant decrease in ultimate lap shear force (ULSF). The Al3Mg1Si/ Ti64 joints were particularly affected and failed at low ULSF of around 1.25 kN , most likely correlated with the lower process temperature due to the lower solidus temperature of the alloy [17].

Joint formation - surface topography.

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Grinding the Ti64 surface using different SiC grinding papers resulted in almost the same ULSF, while varying the Ti64 surface topography by grinding and SB influenced the joint strength significantly. SB led to a drastically reduced ULSF, especially for alloys having only 1wt%Mg. As presented in Table 3, mean Rz and Ra values of the generated surfaces were similar, only the standard deviation was larger for the sand blasted Ti64 surface. However, it is assumed that the effect on the ULSF is correlated rather to the topography than simply roughness characteristics. As compared in Figure 3, in terms of surface structure, the trenches generated by SiC paper grinding differed a lot compared to the peaks and valleys generated by SB . It is likely that within the grinding trenches higher diffusion rates are present. The bonding was enhanced, while the effect of mechanical interlocking is expected to be reduced compared to SB. The effect originated due to the parallel orientation of the grinding trenches to the load direction during tensile testing, see Figure 3d.
The higher diffusion rates may result from Al being continuously plastically deformed within the trenches of the grinded topography, enhancing grain boundary diffusion [8], while the Al being squeezed into the SB structure predominantly staying there for the rest of the welding process.
The samples subjected to griding and knurling showed large variations, where all material combinations, except for the Al1Mg0Si/Ti64 joints, were characterized by high standard deviations. It can be seen by the upper ULSF range that knurling has high potential, but the lower ULSF range indicates that joints are also sensitive to unbonded regions and therefore local crack initiation sites. The results agree with the findings explained in the introduction, which describe the advantages of micro mechanical interlocking, but also the issues correlated with partially unbonded areas [13,14].

Joint formation - alloy chemical composition.

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Based on the content of alloying element, it is shown that the absence of Si , i.e. in AllMg0Si/Ti64 joints, limits the ULSF to roughly 4.5 kN . This can be recognized e.g. in the 1.5 mm PD #320 sample and might be caused by missing Si required for formation of the strengthening precipitates. Additional Si up to 1%(Al1Mg1Si/Ti64) improved the mechanical properties significantly, i.e. the highest ULSF increased to 6.5 kN in PD 1.5 mm #320 sample. However, additional Si more than 1wt%. barely influenced the mechanical properties, as shown in Al1Mg1.5Si/Ti64 sample in 1.5 mm PD #100. Meanwhile, an increase in Mg content, i.e. Al1Mg1Si/Ti64 to Al3Mg1Si/Ti64, led to a lower ULSF in ground samples with 1.5 mm PD in #320 and #120 samples, but resulted in a higher ULSF in for the knurled and SB samples, as shown in Figure 4. These results indicate that a higher Mg content might also play an important role in the mechanical interlocking component of the joint performance.

[picture]

Fig. 4. Ultimate lap-shear forces measured during tensile testing the four Al alloys welded to surface pre-treated Ti64.

Fig. 4. Ultimate lap-shear forces measured during tensile testing the four Al alloys welded to surface pre-treated Ti64.

Fracture behavior.

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Figure 5 presents fracture surfaces of a PD 1.5 mm,#320Al1Mg1.5Si/Ti64 joint. A clear transition between bonded and unbonded region is presented in Figure 5a. It mainly shows that both interfacial fractures as well as through the Al fractures occur at the interface. The fracture sometimes also alternates several times and a clear behavior for correlating the fracture mode with specific interface areas could not be identified, as shown in the interface presented in Figure 6a. When the fracture propagates through the Al , leaving attached Al on the Ti64 side, this Al is characterized by a clear dimple structure indicating a ductile failure behavior, see Figure 5b.

[picture]

Fig. 5. Fracture surface of a) a PD Al1Mg1.5Si/Ti64 Ti64 joint, highlighting the mixed failure mode of partially interfacial delamination and partial through the Al side with Al being attached to the Ti64 surface, and b) a highlighted area of the PD Al/Ti joint fracture surface, showing the fracture locally through the Al in a ductile behavior indicated by the observed dimple structure.

Fig. 5. Fracture surface of a) a PD A l 1 M g 1.5 S i / T i 64 Ti64 joint, highlighting the mixed failure mode of partially interfacial delamination and partial through the Al side with Al being attached to the Ti64 surface, and b) a highlighted area of the PD A l / T i joint fracture surface, showing the fracture locally through the Al in a ductile behavior indicated by the observed dimple structure.

Figure 6a presents the fracture surface of a knurled Al/Ti joint. It is visible that Al is locally attached inside and between the indents, indicating that fracture propagated through the Al, while the fracture surface without adhering Al shows delamination through the mathematical formula interface. Inside the indents, Al is attached to one side of the indents (upper face), while the Al is mainly detached at lower face, Figure 6 a. Consequently, the fracture runs through the interface on the lower side of the indent and then further propagates through the Al. This is based on the differing interlocking
conditions with regard to load direction, as the Al is pulled towards the upper indent face while the Al is pulled away from the lower face. This phenomenon was observed at different locations of the spot weld, including where the rotation direction and loading direction during the tensile test were in the same and opposite direction, respectively. Therefore, the effect is not mainly related to improved bonding conditions on one side of the indent, influenced by the rotation direction of the tool, but mainly due to the load and secondary bending effects during tensile testing. Figure 6b shows a single indent at higher magnification. Although no Al was visible to be attached to the lower indentation face from the macroscale, a small amount of Al was found to be adhered at the microscale. This implies that such surface pre-treatment needs to be handled with care. Individual parameter optimization needs to be performed to enable sufficient diffusion even in the most unsuitable areas for diffusion of the generated surface structure.

[picture]

Fig. 6. SEM micrograph of a knurled superscript number joint fracture surface, showing a) detached Al on one side of the knurling indents. Al is attached on the other side and partially on the Ti64 surface between the indents. Lines of attached Al are visible, indicating discontinuous delamination of the interface, b) Low amount of attached Al fragments within a knurling indent on the mainly detached side.

Fig. 6. SEM micrograph of a knurled joint fracture surface, showing a) detached Al on one side of the knurling indents. Al is attached on the other side and partially on the Ti64 surface between the indents. Lines of attached Al are visible, indicating discontinuous delamination of the interface, b) Low amount of attached Al fragments within a knurling indent on the mainly detached side.

Conclusion

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Refill Friction Stir Spot Welding of AlMgSi with different Mg and Si contents to Ti 6 Al 4 V is investigated for two plunge depth (PD) conditions and various Ti sheet surface pre-treatments. The Mg and Si content was systematically varied, and the influence on the mechanical properties of the obtained spot welds was investigated. The following conclusions are drawn:

The PD plays a significant role in welding dissimilar mathematical formula and a low tool to interface spacing is necessary to obtain sound joints.

Surface pre-treatments have a strong influence on the mechanical properties of the joints. Randomly distributed peaks and valleys generated by sandblasting led to lower mechanical properties compared to grinding trenches with similar roughness values.

Knurling the Ti64 surface to enable additional micro-mechanical interlocking is promising as shown by partially outstanding mechanical properties, but also prone to unbonded regions as can be seen by the high variations of properties.

The fracture surfaces show an alternating crack propagation that led to locally interfacial delamination and fractures through the Al. Attached Al on the Ti64 sheet shows a dimple structure, indicating ductile fracture characteristics.

Acknowledgement

The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

Funding by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - project number 464986536 - is gratefully acknowledged.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.139.pdf

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This ePub file was created using machine-generated methods and automated software to convert the source text into an ePub format and may contain formatting inconsistencies from the original version. Formatting may vary depending on the device used.

The original version of this eBook is available on https://www.scientific.net/book/innovative-joining-by-forming/978-3-0364-1990-9

Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 149-160

doi: 10.4028/p-8jkHA8

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-11

Accepted:
2026-02-23

Online:
2026-04-15

Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

Pia Katharina Kaimann n1,a*, Mathias Bobbert 1, b and Gerson Meschut 1,c

1 Paderborn University, Pohlweg 47-49, 33098 Paderborn, Germany

a  pia.kaimann@lwf.upb.de, b  mathias.bobbert@lwf.upb.de, c  gerson.meschut@lwf.upb.de

(*corresponding author)

Keywords: Self-piercing riveting, FEM, Versatility, Movable die.

Abstract

Self-piercing riveting (SPR) is a well-established joining technique in lightweight construction, as it enables the joining of different materials without requiring pre-drilling. However, the necessary adaptation of the rivet-die combination to the respective material and thickness combinations requires a large number of specific tool sets, which significantly limits the process's flexibility. To overcome these limitations, the versatile self-piercing riveting (V-SPR) was developed, which features enhanced punch actuation in combination with a multi-range-capable rivet [1]. In this context, the concept of a movable die was introduced, which enables an extended process window and adaptable joint formation. Kappe et al. presented initial studies demonstrating the potential of this approach [2]. However, a detailed numerical understanding of the underlying mechanisms remains lacking. This paper presents a numerical analysis of V-SPR with a movable die using a finite element (FE) model. The model includes deformable rivets, sheet metal materials and a kinematically controlled die with adjustable movement. A parameter study was conducted to analyse the influence of die movement on the material flow of the rivet and sheets, as well as joint formation. The simulations were validated using selected experimental data. The goal is to compare the joint geometries achieved with fixed and moving dies and expand the process windows of V-SPR. The results demonstrate that the movable-die concept significantly enhances the material flow of both the sheets and the rivet, resulting in a noticeably larger and more reliable interlock than what is achievable with V-SPR using a fixed die. The numerical analyses support the observations reported by Kappe et al. and extend them by providing a quantitative description of how die displacement influences the resulting interlock size. Moreover, the ability to precisely control the die movement makes it possible to join challenging sheet-metal combinations that are difficult to process with conventional setups, particularly in cases involving thicker sheet materials.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

Reducing CO2 emissions has become a key objective in the automotive industry due to stricter environmental regulations and the EU's target of significantly cutting emissions by 2030 [3]. Alongside new drive concepts, lightweight construction remains one of the most effective strategies, as lowering vehicle mass directly reduces emissions [4]. Modern lightweight design increasingly relies on multi-material systems that combine materials with different properties to meet local load requirements. However, this material mix creates significant challenges for thermal joining methods, which often fail when joining dissimilar materials such as aluminium and high-strength steel due to metallurgical incompatibilities. For this reason, mechanical joining technologies, especially selfpiercing riveting (SPR), have become essential, offering a reliable solution for joining diverse material combinations in automotive applications.

Self-piercing riveting (SPR) enables joining multiple layers, even of different materials, without pre-drilling, provided access is available from both sides. The process sequence is shown in Fig. 1. After positioning the components, the blank holder fixes the sheets. The rivet is then pressed into the die-side layer by the punch, cuts through it and takes up the punched-out remnant in the rivet shank. When the tool-side position is reached, the rivet is radially expanded and compressed by the die contour, creating a form-fitting and force-fitting closing head. After the return stroke of the punch
and blank holder, the connection remains fully formed. The process takes place in a single stroke, can be automated and is very reliable [5].

[picture]

Fig. 1. Process sequence of the self-piercing riveting [5]

Fig. 1. Process sequence of the self-piercing riveting [5]

One restriction of SPR is the need to adapt the rivet-die combination to changing boundary conditions, such as variations in sheet thickness or material. This limitation has driven research toward more versatile and adaptable joining processes that enable broader applications. One such development is versatile self-pierce riveting (V-SPR), which combines extended tool kinematics with a multi-range capable rivet to join varying sheet thickness combinations without altering the rivet-die setup [1]. In contrast to conventional SPR, the punch-side tool is split into an inner and an outer punch. The process sequence is shown in Fig. 2. The inner punch first sets the rivet, after which the outer punch forms the rivet head to the respective sheet thickness, enabling robust joining even under fluctuating joining conditions.

[picture]

Fig. 2. Process sequence of the versatile self-piercing riveting [1]

Fig. 2. Process sequence of the versatile self-piercing riveting [1]

Another approach to reducing rivet-die combinations is the use of a two-part die principle with a movable die element, as shown in [6]. Drossel and Jäckel (2014) were able to show that brittle aluminium material behaviour can be joined by movement in the bottom of the die [7]. Kappe et al. show that both a continuously lowering die bottom and subsequent embossing significantly increase the formation of the closing head. At the same time, material flow and stress states change, with a continuously lowering die bottom in particular generating increased compressive stresses in the slug and thus supporting rivet expansion. Subsequent stamping also improves the connection quality, but places a strain on the tools [2].

This study combines the V-SPR and the concept of a continuously descending die and numerically analyses the influence of different die movement sequences on joint formation, as shown in Fig. 3.

[picture]

Fig. 3. Process sequence of the V-SPR with movable die

Fig. 3. Process sequence of the V-SPR with movable die

Experimental Procedure

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

The rivet, designed for use across multiple material ranges, is made from C35 B+Cr steel. The sheets are made of aluminium alloy EN AW 6014 in temper T4. The chemical compositions of both the rivet material and the EN AW 6014 alloy used in this study are listed in Table 1.

Table 1. Chemical composition properties of rivet material C35 B+Cr [8] and sheet materials EN AW-6014 temper T4 [9]

Chemical composition C35 B+Cr (weight %)
Elements | Si | C | Cu | Mn | P | S | Cr | Mo | B
Min. |  | 0.35 |  | 0.60 |  |  |  |  | 0.0008
Max. | 0.3 | 0.40 | 0.25 | 0.90 | 0.025 | 0.025 | 0.3 | - | 0.005
Chemical composition EN AW-6014, T4 (weight %)
Elements | Si | Fe | Cu | Mn | Mg | Cr | Zn | Ti | V
Min. | 0.30 |  |  | 0.05 | 0.40 |  |  |  |
Max. | 0.60 | 0.35 | 0.25 | 0.20 | 0.80 | 0.20 | 0.10 | 0.10 | 0.10

For the joining operation, a setup is used that enables the inner and outer punches to move independently, as outlined earlier. As illustrated in Fig. 4, the upper tooling, consisting of both punch elements and the blank holder, and the lower tooling are mounted within a column-guided frame [10].

[picture]

Fig. 4. Joining system with extended punch-sided tool actuator technology to process the multi-range capable self-piercing rivets [10]

Fig. 4. Joining system with extended punch-sided tool actuator technology to process the multi-range capable self-piercing rivets [10]

Numerical simulation

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

To examine how a movable die affects joint formation in versatile self-piercing riveting, a twodimensional axisymmetric simulation model was developed, as illustrated in Fig. 4. The approach builds on earlier work by KAPPE et al. [11]. The finite-element model was generated using the LSDyna simulation software and evaluated using an implicit solution strategy. The inner and outer punches were represented as elastic bodies to enable a detailed assessment of the forces generated during the operation. A displacement-controlled punch motion was imposed by constraining the corresponding nodes in the model. The inner punch moves at 9 mm/s. The movable die consists of a rigid die ring and a movable die base, both modelled as elastic parts, with their joining force measured in the lower area. The blank holder, implemented as a rigid body, applied a constant load throughout the entire procedure. Material behaviour was introduced using experimentally derived flow curves for the sheet and rivet material. Contact interactions between the rivet and sheets, as well as between the sheets themselves, were described using a static friction coefficient of 0.2 [12]. The punch-side sheet was cut according to a geometric criterion triggered by the prevailing process conditions.

[picture]

Fig. 5. Numerical setup of the V-SPR joining process

Fig. 5. Numerical setup of the V-SPR joining process

The numerical simulation was validated using the material thickness combination EN AW-6014 in 1.5 mm on the punch side and EN AW-6014 in 2.0 mm on the die side. Fig. 6 compares experimentally produced joints with corresponding simulation results, both in micrographs and in quality-relevant parameters such as interlock formation and the minimum die-sided material thickness. The visual comparison shows that the simulation captures the characteristic deformation behaviour of the sheets and the rivet, including the radial expansion of the rivet shank, the material flow into the interlock zone, and the thinning of the die-side sheet, with a high degree of fidelity. Likewise, the numerical predictions of the quality parameters closely align with the measured values, with only minor deviations. This agreement confirms that the model reliably represents the dominant physical mechanisms of the joining process and can therefore be used with confidence to analyse parameter influences and to explore process optimisation strategies.

[picture]

Fig. 6. Validation of material combination EN AW-6014 in 1.5 mm in EN AW-6014 in 2.0 mm

Fig. 6. Validation of material combination EN AW-6014 in 1.5 mm in EN AW-6014 in 2.0 mm

Results and Discussion

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

To determine the influence of die movement on material flow and joint formation, a full factorial design of experiments (DoE) was set up (see Table 2). The following factors are varied: the thickness of the punch-side sheet metal, the thickness of the die-side sheet metal, the preset die depth at the
start time, the die velocity, and the start time of the die movement, depending on the total process duration. The quality-relevant parameters, interlock (f) and the minimal die-side material thickness (tr), were used as the response variables. The range of the start time of the die movement is based on the total time of the inner punch movement. The percentage increments are selected around the cutting process of the reference connections. The die velocity is set below the inner punch velocity so the die bottom can apply counterpressure.

Table 2. Varied factors of the DoE to determine the influence of the die movement

Factor | Unit | Range
Punch-side sheet | mm | 1.5 | 1.75 | 2.0
Die-side sheet | mm | 1.5 | 1.75 | 2.0
Start time of die movement | % | 35 | 38 | 41
Die velocity | mm/s | 6.5 | 7 | 7.5
Die depth at start | mm | 0 | -0.3 | -0.6

Fig. 7 shows the main-effect diagram of the investigated parameters regarding the influence of die movement on the interlock f and the minimal die-side material thickness tr.

[picture]

Fig. 7. Main effect diagram of the die movement parameter for the interlock t1=1.5 mm and the minimal dieside material thickness t2=1.5 mm, b related to the five parameters investigated: punch side sheet, die side sheet, start time of die movement, die velocity and die depth at start.

Fig. 7. Main effect diagram of the die movement parameter for the interlock t 1 = 1.5 m m and the minimal dieside material thickness t 2 = 1.5 m m , b related to the five parameters investigated: punch side sheet, die side sheet, start time of die movement, die velocity and die depth at start.

For the interlock, it can be seen that the sheet thickness on the punch side, the die depth at the start of the process and the timing of the die start in particular cause significant changes in the interlock. Increasing punch-side sheet thickness significantly reduces interlock (negative correlation), demonstrating that the movable die mechanism becomes critical for thick-sheet applications where conventional SPR typically fails. In contrast, the die depth at the start time has a strong positive effect on interlock formation. The lower the die depth at start time, the earlier the slug fills the rivet shank, cutting the sheet metal on the punch side earlier and increasing interlock, as shown in Fig. 8.

[picture]

Fig. 8. Example simulations for rivet shank filling depending on the die depth at the start time, with the other factors remaining constant. a) die depth at start t1=1.5 mm, b) die depth at start t2=2.0 mm and c) die depth at start t1=2.0 mm

Fig. 8. Example simulations for rivet shank filling depending on the die depth at the start time, with the other factors remaining constant. a) die depth at start t 1 = 1.5 m m , b) die depth at start t 2 = 2.0 m m and c) die depth at start t 1 = 2.0 m m

The start time of the die movement also has a positive influence. The later the die movement begins, the greater the interlock achievable. The die-side sheet thickness and die velocity have no significant effect on interlock formation. The die-side sheet thickness has a slight positive effect. The
die velocity exhibits non-monotonic behaviour, with a medium value enabling the highest interlock, while slower and faster speeds produce lower values.

For the minimal die-side material thickness, the effect profile is partially opposite. Both sheet thicknesses, on the punch and die sides, show a significantly strong positive effect, as greater initial thicknesses favour higher minimum die-sided material thickness. In contrast to the interlock, however, a low die depth at the start of the process has a negative effect, as the rivet penetrates deeper into the material, thereby reducing the minimum die-side material thickness. The start time of the die movement shows a slight negative trend, with later movement times resulting in slightly greater thinning. The die speed has no significant effect on the residual bottom thickness. Overall, it is clear that both the die depth at the start of the process and the temporal and kinematic behaviour of the die movement are decisive for how material flow, radial expansion and thinning interact during the setting and expansion process.

For the material combination EN AW-6014 with sheet thickness a) t2=2.0 mm in t1=1.5 mm ) t2=1.5 mm in t1=1.5 mm, and c) t2=2.0 mm in t1=2.0 mm, the optimisation of the die movement (red) to maximise the interlock to be achieved compared to the refence die (black), with a depth of 1.2 mm , is shown in Fig. 9.

Fig. 9. Joining formation of the optimisation of the movable die (red) and the reference die with a deep of 1.2 mm (black) for the material combination EN AW-6014 with sheet thickness a) t2=2.0 mm in 27.6%, b) 1.5/1.5 mm in 51.8%, and c) 1.5/2.0 mm in 2.0/2.0 mm related to the interlock f

This shows that the die actuator enables early cutting of the punch-side sheet metal. As a result, the punch-side sheet metal is drawn in only slightly, and the die-side sheet metal forms more closely around the rivet shank, thereby increasing the interlock. However, the residual bottom thickness is minimised.

Table 3 illustrates the additional interlocking effect achievable with a movable die. For all sheet thickness combinations examined, the interlock increases significantly compared to the rigid reference die with a 1.2 mm depth. The gain is particularly pronounced at higher package thicknesses: while an increase of around 49.1% is already achieved with the thinnest combination ( 1.5/1.5 mm ), the interlock increases by more than 40.6% with the 1.5/2.0 mm configuration. A similarly strong increase is observed for the thickest combination ( 18.2% ), which still achieves a substantial (2.0/2.0 mm) enhancement. This clearly shows that the active die movement significantly improves the material flow during the setting process and unlocks considerable optimisation potential, especially with thicker sheet combinations.

Table 3. Optimisation potential of the interlock achieved by the active die actuator compared to reference joints

| Combination | Interlock f, die depth 1.2 mm | Interlock f, movable die | Procent
a | EN AW-6014, t₁ = 1.5 mm, t₂ = 1.5 mm | 0.2843 mm | 0.3628 mm | + 27.6 %
b | EN AW-6014, t₁ = 1.5 mm, t₂ = 2.0 mm | 0.2987 mm | 0.4535 mm | + 51.8 %
c | EN AW-6014, t₁ = 2.0 mm, t₂ = 2.0 mm | 0.1398 mm | 0.2085 mm | + 49.1 %

Table 4 shows the resulting reduction in the minimum die-side material thickness as a result of interlock optimisation through the use of an active movable die. For all sheet thickness combinations examined, a significant decrease in the minimum die-sided material thickness is observed compared to the reference with a rigid die and a depth of 1.2 mm . This effect is most pronounced in the thinnest combination ( 9.8% ), where the minimum die-side material thickness is reduced by around 40.6%. For the 1.5/1.5 mm combination, the reduction is around tr, while for the thickest combination 0.2·, it is comparatively moderate at around t1=1.5 mm,t2=1.5 mm. This clearly shows that the increase in interlock achieved by active die movement is accompanied by greater thinning of the dieside sheet metal, with this trade-off particularly pronounced in thinner sheet material combinations. The )t1=1.5 mm reduction in minimum die-side thickness for the t2=2.0 mm combination approaches failure-critical levels. According to DVS 3410 guidelines, t1=2.0 mm,t2=2.0 mm should typically exceed t1=1.5 mm for reliable joint strength [5].

Table 4. Resulting reduction in minimum die-sided material thickness through optimisation of the interlock by active die actuators

Combination

Minimum

die-side material

thickness mathematical formula,

die depth 1.2 mm

Minimum

die-side material

thickness mathematical formula,

movable die

Procent

a | EN AW-6014, superscript reference | 0.1449 mm | 0.0861 mm | mathematical formula
b | EN AW-6014, superscript reference | 0.3521 mm | 0.2881 mm | superscript reference
c | EN AW-6014, superscript reference | 0.6957 mm | 0.6278 mm | mathematical formula

Fig. 10 shows the force-time diagram of the die bottom during the inner punch movement of the optimised die movements. At the beginning, the force increases at a constant die depth of zero for the three material combinations to an approximately equal level of between 18.7 kN and 19.3 kN . For the material thickness combination, a) t2=1.5 mm, the force drops to approximately 8.9 kN when the die movement starts and is kept constant until the final die depth of 1.2 mm is reached. Similar force levels result for the material thickness combination b20% in t1=2.0 mm. When the die movement starts, the force drops to approximately 8.3 kN , which remains constant until the end of the die movement. As with the material thickness combination c) t2=2.0 mm, the force here drops to an almost constant level of 7.6 kN . Once the maximum die depth of 1.2 mm is reached, there is a sharp increase in the force on the die bottom. The maximum force, which is between 33 and 36 kN , is reached at the maximum inner punch stroke. The evaluation of the resultant forces at the die bottom shows that the cutting process of the punchside sheet metal has no effect on the resultant forces.

[picture]

Fig. 10. Force-time diagram of the die base during movement of the inner punch for the material combination EN AW-6014 with sheet thickness a) 20% in t1=1.5 mm, b) t2=1.5 mm in t1=2.0 mm/t2=2.0 mm, and c) 12c) in 27.6%

Fig. 10. Force-time diagram of the die base during movement of the inner punch for the material combination EN AW-6014 with sheet thickness a) 20 % in t 1 = 1.5 m m , b) t 2 = 1.5 m m in t 1 = 2.0 m m / t 2 = 2.0 m m , and c) 12 c ) in 27.6 %

In this strategy, the movement is initiated once the predefined force threshold is reached. The control system then maintains this force by lowering the movable die until either the maximum die depth is attained or the process kinematics prevent further die travel. For the material combination EN AW-6014 with sheet thickness a) 52% in (2.0/2.0 mm), forces of 18.72 kN and a force 9.8% higher at 22.46 kN were selected. For the material combination EN AW-6014 with sheet thickness c) 40.6% in 1.5/1.5 mm, forces of 19.27 kN and a force t1=2.0mm/t2=2.0 mm higher at 23.12 kN were selected. The joint formed by force-controlled die movement in green, which aims to maximise the interlock, compared with the reference die in black with a depth of 1.2 mm , is shown in Fig. 11.

[picture]

Fig. 11. Joining formation of the force-controlled movable die (blue mathematical formula and green superscript number ) and the reference die with a deep of 1.2 mm (black) for the material combination EN AW 6014 with sheet thickness a) superscript number in superscript number and c ) superscript number in superscript number related to the interlock f

Fig. 11. Joining formation of the force-controlled movable die (blue and green ) and the reference die with a deep of 1.2 mm (black) for the material combination EN AW 6014 with sheet thickness a) in and c ) in related to the interlock f

Fig. 12 presents the force-time response of the force-controlled movable die during the inner punch stroke. Initially, the force increases at zero die penetration until the respective setpoint is reached. The controlled force levels correspond to 18.72 kN (a) and 22.46 kN (b) for superscript number / superscript number, and 19.27 kN (c) and 23.12 kN (d) for superscript number. After the onset of the die motion, the force is approximately constant until the end of the die movement. The cutting operation on the punch-side sheet does not result in a noticeable change in the force response. For cases (a) and (b), the maximum die depth is reached, which results in a pronounced force increase during the subsequent inner punch stroke. In contrast, for (c) and (d), the movable die does not reach the maximum die depth due to the intensified radial material flow.

Force control positively affects the formation of the interlock. As the set force increases, the interlock form increases. As shown in Fig. 11 c), a significantly increased radial material flow results in a lower force acting on the die base, meaning that the maximum die depth is not reached, see Fig. superscript number and d). Overall, the pronounced radial material flow has a favourable effect on the required die depth and the maximum joining force, but can potentially negatively impact adjacent joining points due to the sheet metal's radial expansion on the die side.

[picture]

Fig. 12. Force-time diagram of the force-controlled movable die during movement of the inner punch for the material combination EN AW-6014 with sheet thickness a) superscript number in superscript number with forces of 18.72 kN b ) superscript number in superscript number with forces of superscript reference ) superscript number in superscript number with forces of 19.27 kN and d ) superscript number in superscript number with forces of 23.12 kN

Fig. 12. Force-time diagram of the force-controlled movable die during movement of the inner punch for the material combination EN AW-6014 with sheet thickness a) in with forces of 18.72 kN b ) in with forces of ) in with forces of 19.27 kN and d ) in with forces of 23.12 kN

Summary

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

This study presents a numerical investigation of versatile self-piercing riveting (V-SPR) with a movable die, aiming to expand the process window and improve joint formation compared to conventional setups with rigid dies. A two-dimensional axisymmetric finite-element model was developed in LS-Dyna, incorporating deformable rivets and sheets as well as a kinematically controlled die base. The model was experimentally validated using EN AW-6014 sheet combinations, demonstrating high agreement in deformation behaviour and quality-relevant parameters, such as interlock and minimum die-side thickness.

A full-factorial parameter study examined the influence of punch-side and die-side sheet thickness, die depth at process start, die velocity, and the start time of die movement on joint formation. The results show that interlock formation is strongly governed by the punch-side sheet thickness, the initial die depth, and the die start time, while die velocity and die-side thickness have only minor effects. For the remaining die-side material thickness, both sheet thicknesses exert the largest influence, whereas a low die depth at the beginning of the process promotes increased thinning. Comparative simulations between a fixed die and a movable die demonstrate substantial optimisation potential. Depending on the sheet thickness combination, the interlock increases between superscript number and nearly superscript number when using an active die movement. However, this improvement is directly associated with a reduction in the minimum die-side material thickness. Compared to the rigid reference die, the minimum die-side thickness decreases significantly across all investigated combinations, with the greatest reduction observed in thin-sheet packages. While the thickest combination superscript number shows only a moderate decrease of around mathematical formula, the thinning becomes critical for thinner configurations, reaching reductions of up to mathematical formula for the superscript number combination. This highlights a clear trade-off between enhanced interlock formation and increased die-side sheet thinning, which is particularly pronounced for thin sheet materials. The evaluation of the forces on the die bottom during the adjustment process also shows that the characteristic force development is largely independent of the sheet thickness combination. In all configurations, the force on the die base increases to approximately 18 to 19 kN before the die movement begins. As soon as the die begins to move, the force drops to an almost constant level, between 7.6 kN and 8.9 kN , depending on the material combination, and remains stable until the final die depth is reached. Only after the maximum tool displacement has been reached does the force increase sharply again, reaching a peak value
between 33 kN and 36 kN at full inner punch stroke. It is noteworthy that cutting the sheet metal on the punch side has no measurable influence on the resulting tool forces. Building on the force-time response shown in Fig. 10, a force-controlled die-movement strategy was investigated. Here, the motion is initiated once a defined force threshold is reached, and the controller maintains this force by lowering the die base until either the maximum die depth is achieved or the process kinematics limit further die travel. Increasing the set force increases the interlock, while the force response remains largely unaffected by cutting the punch-side sheet. For the thicker sheet package ( superscript number ), intensified radial material flow reduces the force acting on the die base, preventing the maximum die depth from being reached, thereby reducing both the required die depth and the maximum joining force. At the same time, the pronounced radial expansion of the die-side sheet may influence adjacent joining points.

The results confirm that a controlled die displacement significantly enhances material flow and rivet expansion, particularly for thicker sheet packages, enabling more robust and versatile SPR joint formation. Building on these results, future work should focus on experimentally validating the identified parameter interactions across a broader range of material combinations and sheet thicknesses.

Acknowledgement

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

Funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - TRR 285/2 - Project-ID 418701707. The authors thank the German Research Foundation for their organisational and financial support. Data regarding the contents of the publication can be requested at www.trr285.de.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.149.pdf

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F. Kappe, C. Zirngibl, B. Schleich, M. Bobbert, S. Wartzack, and G. Meschut, "Determining the influence of different process parameters on the versatile self-piercing riveting process using numerical methods," Journal of Manufacturing Processes, vol. 84, pp. 1438-1448, 2022, doi: 10.1016/j.jmapro.2022.11.019.

Self-piercing riveting: Properties, processing and applications (Woodhead publishing series in welding and other joining technologies 82). Oxford: WP Woodhead Publ, 2014

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 161-169

doi: 10.4028/p-0TiIhI

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-01-19

Revised:
2026-02-17

Accepted:
2026-02-24

Online:
2026-04-15

A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

The original version of this paper is available on https://www.scientific.net/MSF.1185.161.pdf

Devulapally Deekshith Reddy 1,a*, Tröster Thomas 2, b

¹Paderborn University, Faculty of Mechanical Engineering, Chair of Automotive Lightweight

Design, Warburger Str. 100, 33098 Paderborn, Germany

a* deekshith.reddy.devulapally@uni-paderborn.de, b  thomas.troester@uni-paderborn.de

Keywords: clinched joints, non-rotationally symmetric joints, joining.

Abstract

Clinched joints with non-rotationally symmetric geometries exhibit orientation-dependent mechanical behavior that is commonly neglected in structural-scale simulations. Reuleaux triangle shaped clinched joints, in particular, show pronounced in-plane anisotropy depending on their orientation. While such effects have been studied at joint and specimen scale, their relevance at the structural level remains largely unexplored. In this work, the influence of joint orientation on the bending response of a joined structure is investigated using numerical simulations. A simplified joint replacement model based on the *CONSTRAINED_SPR2 point-connector formulation in LS-DYNA is employed, with parameters calibrated from previously obtained experimental force displacement data. A hat shaped profile structure subjected to three-point bending is analyzed in a parametric study considering variations in joint orientation, joint spacing, and profile geometry. The results show that joint orientation has little influence during the initial deformation phase but becomes increasingly significant at larger displacements, where joint behavior governs load transfer. Orientation dependent effects are found to influence the global force displacement response and local load redistribution among joints, with magnitudes comparable to those induced by changes in joint spacing and structural geometry. The findings confirm that joint orientation effects remain relevant at the structural level and should be considered in the design of structures assembled using non-rotationally symmetric clinched joints.

Introduction

The original version of this paper is available on https://www.scientific.net/MSF.1185.161.pdf

Mechanical joining technologies play a key role in lightweight design, particularly in automotive and transportation applications where thin sheet metals and dissimilar materials must be joined efficiently, reliably, and at low cost[1], [2]. Among these technologies, clinching has gained widespread industrial acceptance due to its ability to join sheets without additional fasteners or thermal input, while maintaining good mechanical performance and process robustness[3]. Conventional clinch joints are typically rotationally symmetric, which results in largely isotropic mechanical behavior under in-plane loading conditions.

Recent developments in clinching technology, however, have led to the introduction of nonrotationally symmetric clinch joints[4], [5], [6]. Unlike conventional joints, these advanced geometries exhibit direction-dependent load-bearing characteristics, particularly under shear loading. Experimental studies at the specimen level have demonstrated that such joints can show significantly different stiffness and strength depending on their orientation relative to the applied load[7]. Moreover, investigations on specimens containing multiple non-rotationally symmetric joints have revealed that joint orientation and rotation can influence the overall load-bearing capacity and deformation behavior of the structure[8].

Despite these advances, a detailed understanding of how non-rotationally symmetric clinch joints interact with surrounding component geometries is still limited. In real components, joints are embedded within complex structural features such as stiffeners, flanges, and profiles, where load transfer mechanisms are governed not only by joint properties but also by local and global deformation modes of the component. Neglecting this interaction may lead to inaccurate predictions of structural performance and suboptimal joint placement during the design phase.

Against this background, the present study aims to numerically investigate the interaction between a non-rotationally symmetric clinch joint and its surrounding component structure using finite element simulations. A simplified joint modeling approach is employed to efficiently represent the essential mechanical characteristics of the joint while enabling systematic parametric studies. A hat shaped profile joined by such clinch joints is analyzed under three-point bending. The test evaluates the influence of joint orientation on the global load-bearing behavior and deformation response of the component. Furthermore, geometric variations of the hat shaped profile are introduced to assess their effect on the global load bearing response.

The study is expected to show that both joint orientation and component geometry play a significant role in governing global structural behavior as well as local joint response. By investigating these interactions, the work contributes to a better understanding of direction dependent clinch joints at the component level. It also provides valuable insights for the design, orientation, and placement of non-rotationally symmetric clinch joints in real engineering structures.

Methodology

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Clinched joints with non-rotationally symmetric geometries exhibit direction dependent mechanical behavior due to their asymmetric load-bearing characteristics. Reuleaux triangle based clinched joints provide a manufacturable joint geometry that introduces pronounced anisotropic in-plane properties while remaining compatible with conventional clinching tools. Previous experimental investigations demonstrated that the asymmetric distribution of neck thickness and interlock around the joint circumference leads to orientation dependent stiffness and energy absorption behavior under shear and bending-dominated loading[7]. These characteristics form the basis for the numerical modelling approach adopted in the present study. The orientation dependent behavior of the Reuleaux triangular clinched joint was first examined at the joint and specimen scale in prior studies[7], [8]. These studies demonstrated that the orientation of individual joints influences load sharing between joints, deformation patterns of the connected sheets, and stress redistribution in the surrounding material. Even for specimens with a limited number of joints, joint orientation effects were shown to affect the global mechanical response. These specimen scale findings provide motivation for the present work. It extends the investigation to a structural scale to assess whether joint orientation effects remain relevant in larger, bending dominated assemblies. (see Fig.1).

[picture]

Fig. 1. Schematic overview of the investigation workflow from single joint to structural component.

Fig. 1. Schematic overview of the investigation workflow from single joint to structural component.

The present study is exclusively numerical. No experimental testing is conducted within this work. Instead, experimentally derived joint characteristics from previous investigations are used as input data for the numerical joint model. The objective is to evaluate the influence of joint orientation on the global and local response of a joined structure subjected to bending loads. By adopting a purely numerical approach, the effect of joint orientation can be isolated in a controlled manner while enabling the investigation of multiple structural and joint configuration variants. The investigated structure is inspired by a previously published three-point bending configuration used to study the bending behavior of joined sheet metal components[9]. To ensure numerical efficiency, the original concept was simplified by removing the covering plates for hat profiles while the main hat profiles are not used as it is. The structure consists of two hat shaped sheet metal profiles connected by multiple clinched joints located at predefined positions. A three-point bending load case is applied by supporting the structure at two outer points and imposing a prescribed vertical displacement at a
central loading point. High-fidelity finite element models of clinched joints based on solid elements can accurately capture local stress states and failure mechanisms, but their high computational cost limits their applicability in large-scale structural simulations involving multiple joints. To enable efficient structural level analyses, a simplified joint replacement modelling strategy is adopted in this study. The joint is represented using the *CONSTRAINED_SPR2 formulation in LS-DYNA, which models the connection as point connector between shell elements, as schematically illustrated in Fig. 2. The connector parameters are calibrated using experimental force displacement data obtained from previous joint and specimen scale investigations. The connector parameters are calibrated using experimentally obtained normal and shear force displacement responses from joint- and specimenlevel tests, while other loading modes are not considered in the calibration procedure. In structural bending simulations, however, the joints may experience combined deformation states. Consequently, although the calibrated connector reliably reproduces force transfer trends under dominant normal and shear loading, quantitative accuracy outside these calibrated loading conditions cannot be strictly guaranteed. The resultant force response of the connector as a function of relative displacement reproduces the experimentally observed joint behavior for specific loading directions and joint orientations. The *CONSTRAINED_SPR2 connector formulation computes joint forces based on elastic-plastic force displacement relations with optional softening behavior. However, the formulation does not explicitly represent physical fracture mechanisms of the joint. The model is intended to capture the global mechanical contribution of the joint with minimal computational cost, enabling its application in structural simulations.

[picture]

Fig. 2. Point-connector representation of the clinched joint using CONSTRAINED_SPR2.

Fig. 2. Point-connector representation of the clinched joint using CONSTRAINED_SPR2.

A parametric study is conducted to investigate the influence of joint orientation and structural geometry on the bending response of the structure. A total of ten clinched joints are distributed equidistantly along the lower hat profile. The center-to-center joint spacing, denoted as 𝐝1, is varied between 35 mm and 40 mm , while the height of the lower hat profile, denoted as 𝐛2, is varied between 11 mm and 22 mm . In addition to these global geometric parameters, the joint orientation is treated as a local parameter, with all joints aligned either at 0∘ or at 180∘ with respect to the in-plane reference direction. The structure is subjected to a three-point bending load case. The structural geometry, joint locations, and parameter definitions ( 𝐝1,𝐛2 ) are shown in Fig. 3.

[picture]

Fig. 3. Geometry and dimensions of the structural model used for the three-point bending simulations (dimensions in mm ) and the joint orientations.

Fig. 3. Geometry and dimensions of the structural model used for the three-point bending simulations (dimensions in mm ) and the joint orientations.

Finite Element Modelling

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The numerical simulations are performed using the explicit solver of LS-DYNA (lsdyna_smp_d_R14.0). The finite element model of the investigated structure, including the three-point bending setup, joint locations, supports, and impactor, is shown in Fig. 4. The impactor is positioned with a lateral offset of 5 mm from the geometric center of the structure. The structural components are discretized using fully integrated shell elements (ELFORM 16) with a constant shell thickness of 2 mm and five through thickness integration points. A nominal element size of 5 mm is used, corresponding to a standard mesh resolution commonly applied in full vehicle simulations. Locally refined meshes are employed in regions with geometric bends to ensure adequate mesh quality and deformation representation. The selected mesh resolution represents a compromise between numerical accuracy and computational efficiency and is consistent with mesh densities commonly used in full vehicle simulations. The sheet material is modelled using a Tabulated Johnson Cook material formulation with an experimentally derived plasticity curve. The material corresponds to aluminium alloy EN AW-6014-T4 and is applied consistently to all deformable sheet components. The supports and impactor are modelled as rigid bodies to avoid local deformation effects. Clinched joints are represented using the *CONSTRAINED_SPR2 point connector formulation. Identical connector parameters are assigned to all joints, while two joint orientation variants ( 0∘ and 180∘ ) are implemented using separate *CONSTRAINED_SPR2 definitions corresponding to the respective joint configurations, without modifying the surrounding mesh or boundary conditions. The connector parameters are obtained from a calibration procedure based on experimental force displacement data from joint and specimen scale tests. A comparison between the experimental responses and the calibrated connector behavior is shown in Fig. 5. The interactions between sheet components are modelled using an automatic surface-to-surface contact formulation. Sheet-to-sheet contact is activated throughout the simulations to prevent interpenetration during deformation. The structure is subjected to a three-point bending load case. The supports are modelled as rigid bodies constrained in all translational and rotational degrees of freedom. The impactor is modelled as a prescribed rigid body with translational freedom only in the loading direction. A prescribed displacement corresponding to a total impactor travel of 30 mm is applied over a simulation time of 1.2 s . The simulations are conducted without mass scaling, and the loading rate is selected to ensure stable numerical behavior. Although a tabulated Johnson-Cook material formulation is used, strain-rate effects are not explicitly evaluated. The global structural response is evaluated using the force displacement behavior at the impactor, while local response quantities are obtained by extracting the resultant forces transmitted through the individual *CONSTRAINED_SPR2 joint connectors.

[picture]

Fig. 4. Finite element model setup of the structure and three-point bending configuration.

Fig. 4. Finite element model setup of the structure and three-point bending configuration.

[picture]

Fig. 5. Calibration of the *CONSTRAINED_SPR2 joint model based on experimental forcedisplacement data.

Fig. 5. Calibration of the *CONSTRAINED_SPR2 joint model based on experimental forcedisplacement data.

Results and Discussion

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Global Bending Response of the Structure

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The global bending response of the structure is evaluated using the force-displacement behavior at the impactor, as shown in Fig. 6. The results are compared for different combinations of joint spacing and joint orientation, denoted as D40_J0, D40_J180, D35_J0, and D35_J180, where D35 and D40 represent joint spacings of 35 mm and 40 mm , respectively, while J0 and J180 denote joint orientations of 0∘ and 180∘. Up to an impactor displacement of approximately 20 mm , all configurations exhibit nearly identical force-displacement behavior. In this displacement range, the force increases monotonically to approximately 6−7kN with only minor differences between variants. This indicates that the response is primarily governed by the global bending stiffness of the structure and contact interactions between components, while the influence of joint orientation and spacing remains negligible. In this regime, the joints mainly act as kinematic constraints, and orientationdependent joint behavior does not yet control the global response. Beyond an impactor displacement of approximately 20 mm , a clear divergence between the curves becomes evident. In this deformation regime, load transfer through the joints becomes progressively more significant, and the structural response increasingly depends on the joint configuration. Differences observed at larger displacements can therefore be attributed primarily to orientation-dependent joint behavior rather than to global structural nonlinearities alone. Both joint orientation and joint spacing influence the transmitted force and the evolution of stiffness degradation. Among the investigated configurations, D35_J180 exhibits the highest load-carrying capacity up to an impactor displacement of
approximately 25 mm , indicating a comparatively stiff global response. However, this variant also shows a pronounced and abrupt force reduction beyond this displacement level. In contrast, the remaining configurations (D35_J0, D40_J0, and D40_J180) exhibit a more gradual reduction in force, with the onset of stiffness degradation occurring earlier. This behavior suggests a trade-off between load-carrying capacity and post-peak stability at the structural level. While the D35_J180 configuration sustains higher forces over a larger displacement range, it also experiences a rapid reduction in force once a critical deformation level is reached, whereas the other configurations display a smoother transition into the post-peak regime, indicating a more progressive redistribution of load within the structure. Overall, the results demonstrate that orientation-dependent joint effects remain latent during the early, component-dominated bending phase but become increasingly significant at larger deformations, where joint behavior governs load transfer. These findings extend earlier specimen-scale observations to the structural level and confirm that the orientation of nonrotationally symmetric clinched joints can influence not only the maximum load capacity but also the stability and evolution of the force displacement response in bending-dominated assemblies.

[picture]

Fig. 6. Global force-displacement response of the structure for different joint configurations. D35 and D40 denote joint spacings of 35 mm and 40 mm , respectively, while J0 and J180 represent joint orientations of 0∘ and 180∘.

Fig. 6. Global force-displacement response of the structure for different joint configurations. D35 and D40 denote joint spacings of 35 mm and 40 mm , respectively, while J0 and J180 represent joint orientations of 0 ∘ and 180 ∘ .

Influence of Lower Hat Profile Height on Global Response

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The influence of the lower hat profile height on the global bending response is evaluated by comparing configurations with profile heights b2=11 mm and b2=22 mm, each investigated for both joint orientations (J0 and J180). The corresponding force-displacement responses are shown in Fig. 7, where the notations b22_J0, b22_J180, b11_J0, and b11_J180 are used. A pronounced influence of the profile height on the global response is observed over the entire displacement range. Increasing the lower hat profile height from b2=11 mm to b2=22 mm leads to a significantly less stiff structural response, as evidenced by the higher force levels attained for comparable displacements. This behavior is consistent with the increased bending stiffness associated with the larger section height and indicates that the global response is strongly governed by structural geometry. Despite this dominant geometric effect, the influence of joint orientation remains clearly visible. For both profile heights, differences between the J0 and J180 configurations emerge beyond the initial deformation phase. Similar to the observations made for the joint spacing study, the force displacement curves for different orientations are nearly identical at small displacements, indicating a component dominated response. With increasing displacement, orientation-dependent differences become more pronounced, demonstrating that joint behavior increasingly contributes to the global response as deformation progresses. For the b2=22 mm configurations, the global response is characterized by a
higher load-carrying capacity but also by a more abrupt loss of stiffness at larger displacements, particularly for the b22_J0 configuration, which exhibits a pronounced force drop. In contrast, the b2=11 mm configurations show lower peak forces but a comparatively smoother post-peak behavior, with a more gradual reduction in force. This indicates that while increasing the profile height changes global stiffness and load capacity, it may also lead to a less progressive redistribution of load once critical deformation levels are reached. The consistent presence of orientation-dependent effects for both profile heights demonstrates that the influence of joint orientation does not vanish when the global stiffness of the structure is significantly altered. Instead, joint orientation and structural geometry interact, affecting not only the magnitude of the global force response but also the stability and evolution of the force displacement behavior in the post peak regime.

[picture]

Fig. 7. Global force-displacement response for different lower hat profile heights and joint orientations. b11 and b22 denote profile heights of 11 mm and 22 mm , respectively, while J0 and J180 represent joint orientations of 0∘ and 180∘.

Fig. 7. Global force-displacement response for different lower hat profile heights and joint orientations. b11 and b22 denote profile heights of 11 mm and 22 mm , respectively, while J0 and J180 represent joint orientations of 0 ∘ and 180 ∘ .

Local Joint Shear Force Response and Load Redistribution

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To investigate the origin of the observed differences in global structural response, the local shear forces transmitted through the clinched joints are analyzed. The evaluation focuses on the first four joints along the lower hat profile, as the remaining joints exhibit similar trends and including all ten joints would not provide additional insight. This selection allows a clear assessment of load redistribution while avoiding redundant information. Fig. 8 shows the shear force displacement response of the first four joints for the 0∘ joint orientation (Fig. 8(a)) and the 180∘ joint orientation (Fig. 8(b)). For both orientations, the joints initially exhibit a similar increase in shear force, indicating a relatively uniform load sharing during the early stages of deformation. This behavior is consistent with the component-dominated response observed in the global force displacement curves. With increasing displacement, a redistribution of shear forces among the joints becomes apparent. In the 0∘ orientation (Fig. 8(a)), the shear forces are more unevenly distributed, with certain joints reaching higher force levels earlier than others. This indicates that the load transfer becomes increasingly localized as deformation progresses, leading to a progressive concentration of shear force in specific joints. In contrast, the 180∘ orientation (Fig. 8(b)) exhibits a more gradual redistribution of shear forces among the analyzed joints. The shear force curves show a more balanced progression, with peak forces occurring at larger displacements and with less abrupt changes compared to the 0∘ configuration. This suggests that the joint orientation influences the manner in which shear loads are shared and redistributed along the joint line. At larger displacements, a reduction in shear force is observed for individual joints in both configurations, corresponding to the degradation of global stiffness seen in the force displacement response. Notably, the more abrupt force drop observed in some global configurations is reflected at the joint level by a rapid decrease in shear force in one or
more joints, whereas configurations with a smoother global response show a more progressive reduction in joint forces. Although the simulations are performed using the explicit solver, the loading rate is selected to approximate quasi-static bending conditions. Quasi-static behavior is verified by monitoring the ratio of kinetic to internal energy throughout the simulation, where the kinetic energy remains negligible compared to the internal energy over the entire loading duration. This indicates that inertial effects are minimal and that the structural response is governed primarily by deformation rather than dynamic effects. Consequently, strain-rate effects are not explicitly evaluated, and the results are interpreted in terms of comparative structural response trends.

[picture]

Fig. 8. Shear force-displacement response of the first four clinched joints for (a) 0∘ joint orientation and (b) 180∘ joint orientation.

Fig. 8. Shear force-displacement response of the first four clinched joints for (a) 0 ∘ joint orientation and (b) 180 ∘ joint orientation.

Overall, the joint-level shear force analysis demonstrates that joint orientation affects not only the magnitude of local forces but also the stability of load sharing among joints. These findings provide a mechanistic explanation for the differences observed in the global bending response and confirm that orientation-dependent joint behavior governs load redistribution at the structural level.

Conclusion

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Clinched joints with non-rotationally symmetric geometries exhibit orientation dependent mechanical behavior that is often neglected in structural simulations. In this work, the influence of joint orientation on the bending response of a joined structure was investigated using a simplified, joint model calibrated from previously obtained experimental data. The numerical results demonstrate that joint orientation has a negligible influence during the initial, component-dominated deformation phase but becomes increasingly significant at larger displacements, where load transfer through the joints governs the global response. Differences in joint orientation led to measurable variations in global force displacement behavior, including changes in load bearing capacity, stiffness degradation, and post-peak stability. These effects are comparable in magnitude to those induced by changes in joint spacing and structural geometry. Local joint force analyses revealed that joint orientation affects load redistribution among individual joints, influencing both the magnitude and the progression of shear forces along the joint line. Configurations exhibiting higher global load capacity were associated with more pronounced load concentration and less progressive redistribution, while other configurations showed a smoother redistribution of joint forces and a more gradual global response. The findings confirm that orientation-dependent joint behavior observed at joint and specimen scale remains relevant at the structural level. Joint orientation can therefore be considered an additional design parameter in structures assembled using non-rotationally symmetric clinched joints. As the present study is based solely on numerical simulations, the results should be interpreted as modellingbased trends rather than quantitatively validated predictions. Experimental verification at the component level is therefore required to confirm the observed structural effects. Nevertheless, the adopted modelling strategy provides a computationally efficient framework for incorporating
orientation-dependent joint behavior into large-scale structural simulations. Future work will focus on component-level experimental validation and on extending the approach to mixed joint orientations and additional loading conditions.

Acknowledgement

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This research was funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation), project number 418701707/2 - TRR 285 Phase 2 (Sub-project B01). The authors gratefully acknowledge this financial support. Data regarding the contents of the publication can be requested at www.trr285.de. They also extend their sincere thanks to the Paderborn Center for Parallel Computing ( PC2 ) for providing the computational resources necessary to carry out this research.

References

The original version of this paper is available on https://www.scientific.net/MSF.1185.161.pdf

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X. Gao, C. Chen, X. Ren, and H. Peng, "Experimental investigation on the mechanical square clinching process for joining A15052 sheets," Proc. Inst. Mech. Eng. B J. Eng. Manuf., vol. 237, no. 5, pp. 791-800, Apr. 2023, doi: 10.1177/09544054221112750;WEBSITE:WEBSITE:SAGE;JOURNAL:JOURNAL:PIBB; WGROUP:STRING:PUBLICATION.

X. Ran, C. Chen, H. Zhang, and Y. Ouyang, "Investigation of the clinching process with rectangle punch," Thin-Walled Structures, vol. 166, p. 108034, Sep. 2021, doi: 10.1016/J.TWS.2021.108034.

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C. Steinfelder, C. Acksteiner, and A. Brosius, "A new joint with versatile properties based on a Reuleaux triangle geometry," CIRP Annals, vol. 74, no. 1, pp. 405-409, Jan. 2025, doi: 10.1016/J.CIRP.2025.03.002.

D. R. Devulapally, C. Steinfelder, T. Tröster, and A. Brosius, "Impact of non-rotationally symmetric joint orientation on neighbouring joints and component performance in lap shear specimens," MATEC Web of Conferences, vol. 408, p. 01035, 2025, doi: 10.1051/MATECCONF/202540801035.

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Materials Science Forum, ISSN: 1662-9752, Vol. 1185, pp 171-183

doi: 10.4028/p-VBNO2g

© 2026 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2026-02-05

Revised:
2026-03-06

Accepted:
2026-03-06

Online:
2026-04-15

Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

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Shengjing Liu 1,a*, Yanhong Mu 1, b, Xiaohua Li 1,c, Hailong Lei 1, d

1 AVIC Manufacturing Technology Institute, Beijing 100024, China

a sjliu_mti@163.com, b muyanhong666@qq.com, c13811485985@163.com,d Ihl_fantasy@126.com 

Keywords: Gas pressure loading curve, Superplastic forming/Diffusion bonding (SPF/DB), Interface weld ratio, Void, Four-layer sandwich structure.

Abstract

This paper proposes, for the first time, that the gas pressure loading path significantly influences the weld interface ratio of SPF/DB four-layer sandwich structures. Based on a gas pressure loading curve incorporating back pressure, experimental verification and analysis were conducted. Ultrasonic C-scanning, metallographic examination, and scanning electron microscopy (SEM) were employed to observe and analyze unwelded defects, elucidating their causes and formation mechanisms. Key process parameters-including back pressure time, back pressure value, and gas inlet delay time-were extracted and defined. The influence of these parameters on the weld interface ratio of four-layer sandwich structures was systematically investigated. Finally, with the objectives of eliminating surface grooves and achieving a high weld interface ratio, reference ranges for these process parameters are provided.

Introduction

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With the advancement of aerospace technology, structural components are increasingly required to exhibit high integrity, specific strength, forming accuracy, temperature resistance, and costeffectiveness. Titanium alloy superplastic forming/diffusion bonding (SPF/DB) technology offers a viable solution to meet these demands.

At elevated temperatures around 900∘C, titanium alloys exhibit low deformation resistance and excellent ductility, enabling simultaneous precision forming against a die and solid-state diffusion bonding under argon gas pressure. The SPF/DB process is characterized by its simplicity, flexibility in stiffener layout design, and capability to integrally form complex, lightweight, and high-strength components with tight dimensional tolerances-making it a representative near-net-shape manufacturing technique.

The SPF/DB process enables the production of hollow multilayer titanium alloy structures that are difficult or impossible to fabricate using conventional methods. To date, various configurations have been successfully developed worldwide, including single-layer components such as frames and caps; double-layer structures such as doors, panels, and covers; three-layer components such as wide-chord fan blades; and four-layer structures including fins, fairings, main landing gear doors, wing access panels, and airfoil components.

In recent years, four-layer sandwich structures manufactured by SPF/DB have found increasing application in aerospace vehicles due to their efficient material utilization and superior load-bearing capacity. However, due to limitations inherent in the forming process and structural design, achieving a high weld interface ratio remains a critical technical challenge that restricts the qualification rate and broader application of these structures. This issue directly affects the loadbearing strength, fracture toughness, and fatigue performance of the final product.

The weld interface ratio in titanium alloy SPF/DB four-layer sandwich structures is influenced by several key factors, including diffusion bonding temperature, pressure, holding time, surface condition, environmental atmosphere, and material microstructure. Numerous studies have systematically investigated these parameters, identifying their individual effects on the weld interface ratio and establishing optimal process windows through experimental validation.

Nevertheless, the influence of the gas pressure loading curve on the weld interface ratio of SPF/DB four-layer structures has not been previously explored. This paper presents, for the first time, the significant role of the gas pressure loading curve in determining the weld interface ratio. Through experimental process optimization, four-layer test specimens with a weld interface ratio of 95% or higher were successfully obtained.

Experimental Materials and Methods

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Experimental Materials

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The experiment utilized TC4 titanium alloy, a medium-strength α+β duplex-phase fine-grained alloy. Due to its high strength, excellent corrosion resistance, good weldability, and biocompatibility, TC4 is widely used in aerospace, medical, and chemical fields. TC4 exhibits good superplasticity and diffusion bonding performance at elevated temperatures ranging from 890∘C to 930∘C.

Microstructural analysis of the as-received material with a thickness of 0.6 mm was conducted. Average values from four test groups are presented in Table 1. The sheet surfaces were bright, flat, and free from scratches or visible pitting.

Table1. Microstructure and Properties of Ti-6Al-4V Sheet

Microstructure State

Microstructure

State Grain Size

Composition (%)

V | Fe | C | N | H | O |  |
Equiaxed duplex | Grade 14 | 5.76 | 4.02 | 0.16 | 0.019 | 0.008 | 0.0035 | 0.128 |

Experimental Methods

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The primary process parameters affecting SPF/DB quality are temperature, pressure, and time. The SPF equipment used in this study features precise temperature control, automatic gas pressure loading control, and mechanical pressure loading systems, along with corresponding automatic detection, recording, and signal feedback systems. This enables accurate control of temperature, gas pressure loading curve, and mechanical pressure.

A typical SPF/DB process for a four-layer sandwich structure is illustrated in Fig. 1. When the mold temperature reaches SPF conditions, argon gas is introduced into the outer layer, forming the outer skin against the die. After a holding period, diffusion bonding is completed at the areas of the inner skin not coated with stop-off agent. The outer layer pressure is then released, and gas is introduced into the inner layer. The inner skin undergoes superplastic forming to create the stiffener structure. Finally, after another holding period, the triangular regions of the stiffeners progressively reduce, and the contact area between the inner and outer skins gradually increases, ultimately completing diffusion bonding both between the inner and outer skins and within the stiffener regions.

From the above analysis, it is evident that the diffusion bonding process between the inner and outer skins is complex. When gas is introduced into the outer layer, a vacuum condition cannot be maintained; otherwise, "full-groove" surface forming defects would appear. The diffusion bonding process occurs concurrently with the stiffener forming process. Shear stress exists between the sheets, characteristic of dynamic diffusion bonding, which tends to result in a relatively poor weld interface ratio and makes the process prone to poor bonding and surface groove defects. In summary, the pressure path effects are particularly critical in four - layer structures.

To improve forming efficiency and the quality stability of stiffener diffusion bonding, a process sequence involving initial core sheet diffusion bonding followed by SPF/DB of the four-layer structure was adopted. The specific process flow is: sheet blanking → degreasing and pickling → pattern preparation → core sheet pocket sealing welding → core sheet diffusion bonding → core sheet cutting → alkali etching and pickling → four-layer pocket sealing welding → four-layer structure SPF/DB. Methods such as degreasing and pickling or alkali etching and pickling are
employed during the process to remove surface oil stains and other impurities, ensuring the cleanliness required for the surfaces to be diffusion bonded.

[picture]

Fig. 1 SPF/DB process for four-layer sandwich structure

Fig. 1. SPF/DB process for four-layer sandwich structure

The four-layer SPF/DB test specimen measures 400 mm in length and 300 mm in width. After superplastic forming, the cavity height between the inner sheets is 8−10 mm.The argon gas purity used for the repeated tests was 99.99%.

Experimental research indicates that during SPF/DB, a gas pressure loading curve based on back pressure can reduce premature local adhesion between the inner and outer skins and effectively suppress surface groove defects in four-layer structures. As shown in Fig. 2, the gas pressure loading curve developed by the SPF/DB team at AVIC Manufacturing Technology Institute, based on years of process development and accumulated experience, effectively suppresses the occurrence of surface grooves.

[picture]

Fig. 2 Gas pressure loading curve based on back pressure

Fig. 2. Gas pressure loading curve based on back pressure

To ensure part forming quality and facilitate research, other process parameters were fixed as follows: SPF/DB temperature at 910∘C±5∘C, maximum inner layer pressure at 21 bar, holding time at 140 min , inner/outer layer pressurization rate at 0.78bar/min, and inner/outer layer depressurization rate at 2.3bar/min.After the back pressure was released, during the holding stage, the outer sheet gas pressure was maintained at 1.0-1.2 bar to prevent surface groove defects.

After SPF/DB completion, ultrasonic C-scanning and metallographic examination were used to analyze the diffusion bonding quality at the four-layer structure interface. Scanning electron microscopy (SEM) was employed for morphological observation and compositional analysis at diffusion bonding defect locations.

Analysis of Experimental Results

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Analysis of Unwelded Defects

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

Currently, the effective technical method for detecting unwelded defects is ultrasonic testing, which automatically calculates the weld ratio based on the ratio of the cumulative area of defects shown in the image to the total area of the part.The core of converting ultrasonic C-scan data into welded/unwelded areas lies in threshold calibration. This process essentially involves binarizing the collected continuous ultrasonic signals (such as amplitude and wave height) by setting one or more critical values, thereby distinguishing between "acceptable" and "defective" areas.

In this experimental procedure, the C-scan was performed using a 50 MHz probe frequency with a 0.5 mm step size. Calibration was performed using a reference specimen with a φ0.8 mm flatbottomed hole, made of the same material as the workpiece to be inspected. After the specimen was certified as qualified by a third-party inspection agency, it was used as the basis for calibration. The probe was focused on the defect location in the standard reference block, and the scanning depth and amplitude were adjusted. The amplitude of the reflected wave generated by the defect was adjusted to approximately 95% of the full screen height, and the defect wave identification threshold was set at 90%. This established the correspondence between defect size and the C-scan image.

With an gas inlet delay time of 5 min , a back pressure value of 13 bar, and a back pressure time of 50 min , typical ultrasonic detection results for interface unwelded defects are shown in Fig. 3. The unwelded defects appear nearly circular and are uniformly and dispersedly distributed along the diffusion interface. However, within 10 mm on either side of the stiffeners, the weld interface ratio is good.

[picture]

Fig. 3 Ultrasonic detection results of interface unwelded defects

Fig. 3. Ultrasonic detection results of interface unwelded defects

Samples from the unwelded area were extracted using high-pressure water jet cutting for supplementary metallographic examination. The metallographic analysis morphology is shown in Fig. 4. The microstructure consists of equiaxed primary α+β, with an α phase content of no less than 50%. No oxide layer or α-rich layer was observed at the diffusion bonding interface or within unwelded areas. The metallographic analysis results show good consistency with the ultrasonic nondestructive testing results. The unwelded defects are approximately circular, with sizes ranging from Φ1μ m to Φ5μ m.

[picture]

Fig. 4 Metallographic examination results of interface unwelded defects

Fig. 4. Metallographic examination results of interface unwelded defects

To further determine the cause of the unwelded defects, additional verification tests were conducted. The diffusion bonding quality at the leading-edge position of the test specimen exhibiting interface unwelded defects was analyzed and tested. The weld interface ratio consistently exceeded 95%, ruling out the influence of material microstructure, temperature, time, and blank surface condition on the dispersive interface unwelded defects. During SPF/DB, the space between the inner and outer skins was maintained under vacuum while keeping other parameters unchanged. The interface diffusion bonding quality of the resulting specimens was analyzed and tested, and their weld interface ratios all exceeded 95%, ruling out the influence of diffusion pressure on the unwelded defects. It is concluded that the defects are closely related to the argon gas present between the inner and outer layers.

To confirm whether the argon gas was contaminated by oxygen or other impurities, SEM observation was performed on typical unwelded defects. A representative morphology of an unwelded defect is shown in Fig. 5. Spectrum analysis was conducted at four locations within the SEM observation area. Locations 1, 2, and 4 correspond to unwelded defect composition analysis, while location 3 represents the average composition analysis of the base microstructure. The spectrum analysis results are shown in Fig. 6, with statistical analysis provided in Table 2. The SEM spectrum analysis results indicate that the main elemental compositions at all four locations are Ti,Al, and V , with no oxygen or other impurity elements detected. The Al content at locations 1 and 2 is higher than at location 4, while the V content is lower. It is inferred that locations 1 and 2 correspond to the α phase, and location 4 corresponds to the β phase. Location 3 represents the average of the base microstructure, with elemental content falling in the intermediate range.

The SEM analysis concludes that the unwelded area exhibits typical void morphology, free from oxides or other impurities.

[picture]

(a) Low-magnification image

Figure 5. (a) Low-magnification image

[picture]

(b) Marking of spectrum analysis locations

Figure 6. (b) Marking of spectrum analysis locations

Fig. 5 SEM images of interface unwelded defects

[picture]

Fig. 6 SEM spectrum analysis results

Fig. 6. SEM spectrum analysis results

Table 2 Composition percentages from SEM spectrum analysis

Location | Ti (at%) | Al (at%) | V (at%)
1 | 89.32 | 6.84 | 3.84
2 | 89.45 | 6.71 | 3.84
3 | 89.87 | 6.29 | 3.84
4 | 89.98 | 5.18 | 4.84

Voids are inevitable intermediate products during the diffusion bonding process and have been extensively studied. The diffusion bonding process generally consists of three stages: (1) Local areas of the interface enter the physical contact stage under plastic and viscoplastic deformation mechanisms; (2) Under diffusion pressure and temperature, atoms become activated, and initial contact areas at the interface begin to form chemical bonds, such as metallic bonds, leading to the appearance of voids; (3) Under surface source and interface source mechanisms, atoms undergo interface and volume diffusion, causing voids to continuously shrink and gradually disappear, ultimately forming a reliable diffusion-bonded interface.

It is inferred that during SPF/DB of the four-layer structure, under the influence of outer layer pressure and the pressure differential between inner and outer layers, the inner and outer skins gradually form against the die and come into local contact. In the second stage of diffusion bonding, approximately dispersed closed voids form at the diffusion interface. Before this stage commences, if the gas from the outer layer cannot be promptly vented, it becomes sealed within these voids. During the third stage, under surface source and interface source mechanisms, atomic interface and volume diffusion occur, causing the voids to shrink continuously. The internal residual gas pressure gradually increases, eventually reaching equilibrium with the force from diffusion mechanisms. At this point, void shrinkage stagnates, resulting in unwelded defects.

The triangular stiffener areas are the last to form. When closed voids form at these locations, the back pressure has already been vented, leaving no residual gas inside the voids. Consequently, within 10 mm on either side of the triangular stiffener area, the interface weld ratio is good.

Influence of Back Pressure Time on Interface Weld Ratio

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

With an gas inlet delay time of 5 min and a back pressure value of 13 bar, back pressure time was set to 20 min , 30 min , 40 min , 50 min , and 60 min , respectively. The ultrasonic detection results for interface weld ratio corresponding to different back pressure time is shown in Fig. 7. The influence of back pressure time on interface weld ratio, derived from these results, is presented in Fig. 8.

Experimental results show that when back pressure time ≤30 min, the interface weld ratio is ≥95%. When back pressure time >30 min, the interface weld ratio gradually decreases with increasing back pressure time, initially with a gentle decline. Between 50 and 60 min , the interface weld ratio drops rapidly from 80% to approximately 50%. Unwelded defects appear as dispersed spots, relatively uniformly distributed along the diffusion bonding interface. Within 10 mm on both sides of the stiffeners, the weld ratio remains good.

[picture]

Fig. 7 Ultrasonic detection results of diffusion bonding quality under different back pressure time conditions

Fig. 7. Ultrasonic detection results of diffusion bonding quality under different back pressure time conditions

[picture]

Fig. 8 Analysis of the influence of back pressure time on interface weld ratio

Fig. 8. Analysis of the influence of back pressure time on interface weld ratio

The metallographic analysis morphologies corresponding to different back pressure time is shown in Fig. 9. No oxide layer or α-rich layer was observed at the diffusion bonding interface or within unwelded areas. The metallographic analysis results show good consistency with the ultrasonic non-destructive testing results. When the back pressure time exceeds 30 minutes, metallographic analysis results show that void-type lack of fusion defects begin to appear at the interface to be diffused, and the length of the void line increases with longer back pressure time. When the weld ratio is high, unwelded defects are nearly circular, with sizes ranging from Φ1μ m to Φ8μ m. When the weld ratio is low, unwelded defects are elliptical or elongated, with the long axis oriented along the diffusion interface, reaching a maximum length of approximately 10μ m.

[picture]

(e) 60 min

Figure 10. (e) 60 min

Fig. 9 Metallographic examination results of diffusion bonding quality under different back pressure time conditions

It is inferred that as back pressure time increases, the difficulty of venting gas from the outer layer gradually increases, leading to a continuous increase in the amount of residual gas trapped within the voids. Consequently, the void volume at shrinkage stagnation becomes larger, and voids do not have sufficient time to evolve into multiple smaller voids, typically appearing elliptical or elongated. Ultimately, the area of unwelded interface defects increases, and the weld ratio deteriorates.

Influence of Back Pressure Value on Interface Weld Ratio

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

With an gas inlet delay time of 5 min , back pressure time was set to 20 min,30 min,40 min,50 min , and 60 min , respectively. For each back pressure time, back pressure values were set to 3 bar, 8 bar, 13 bar, and 18 bar. The influence of different back pressure values on the interface weld ratio was experimentally analyzed for each back pressure time.

Experimental results show that when the back pressure values were 3 bar and 18 bar, respectively, surface grooves of varying sizes appeared at local positions on the parts for different back pressure time. It is inferred that when the back pressure value is low, the lubricating effect is insufficient, leading to premature local adhesion between the inner and outer skins during SPF/DB. As the stiffeners form, the outer skin in the adhered areas is pulled by the inner skin, forming surface grooves. When the back pressure value is high, gas inlet uniformity on the upper and lower sides is compromised, making the inner skin prone to uneven deformation, which can cause unilateral adhesion and result in surface grooves.

The influence of back pressure value on interface weld ratio is shown in Fig. 10. Ultrasonic nondestructive testing and metallographic results indicate that the morphology and distribution of unwelded defects are similar to those previously described. Analysis of experimental results shows that when the back pressure time ≤30 min, under different back pressure conditions, the interface weld ratio remains good, all ≥95%. When the back pressure time exceeds 30 min , under the same back pressure time condition, the interface weld ratio increases with increasing back pressure value, with this effect being particularly pronounced at longer back pressure time.

It is inferred that under the same back pressure time condition, an increase in back pressure value reduces the pressure differential between the inner and outer layers, slowing the deformation rate of the inner skin and delaying the contact time between the inner and outer skins (equivalent to a reduction in effective back pressure time). This reduces the amount of gas sealed within the voids, ultimately improving the interface weld ratio.

[picture]

Fig. 10 Analysis of the influence of back pressure value on interface weld ratio

Fig. 10. Analysis of the influence of back pressure value on interface weld ratio

Influence of Gas Inlet Delay Time on Interface Weld Ratio

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

With a back pressure value of 13 bar, back pressure time was set to 20 min,30 min,40 min,50 min , and 60 min , respectively. For each back pressure time, gas inlet delay time was set to 0 min,5 min , 10 min , and 15 min . The influence of different gas inlet delay time on the interface weld ratio was experimentally analyzed for each back pressure time.

Experimental results show that when the gas inlet delay time was 0 min and 15 min , respectively, surface grooves of varying sizes appeared at local positions on the parts for different back pressure time. It is inferred that when the gas inlet delay time is 0 min (i.e., simultaneous inner and outer layer gas inlet), the gap between the inner and outer skins is small, making local adhesion and subsequent surface groove formation more likely. When the gas inlet delay time is long, the coordinated action of gas inlet into the inner and outer cavities becomes less effective. Under the initial back pressure, the inner skin is prone to uneven deformation, which can cause local adhesion and lead to surface grooves.

The influence of gas inlet delay time on the interface weld ratio is shown in Fig. 11. Ultrasonic non-destructive testing and metallographic results show that the morphology and distribution of unwelded defects are similar to those previously described. Experimental results show that when the back pressure time ≤30 min, under different gas inlet delay time conditions, the interface weld ratio is ≥95%. When the back pressure time exceeds 40 min , the interface weld ratio improves with increasing gas inlet delay time.

It is inferred that when the inner skin initially deforms upon gas inlet, the back pressure increases, the deformation rate decreases, and the contact time between the inner and outer skins is relatively delayed (equivalent to a reduction in effective back pressure time). This reduces the amount of gas sealed within the voids, thus improving the interface weld ratio.

Analysis and Discussion of Other Influencing Factors

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

Factors such as the wall thickness of the inner and outer layers of the four-layer structure, the height of the stiffeners within the cavity, and the design of the inlet/outlet channels are also likely to significantly influence the interface weld ratio and the occurrence of surface grooves, potentially interacting with the parameters investigated in this study.

All results presented in this paper were obtained under fixed conditions for the above-mentioned factors. Due to space limitations, a detailed discussion and analysis of these additional influencing factors are not provided here.

[picture]

Fig. 11 Analysis of the influence of gas inlet delay time on interface weld ratio

Fig. 11. Analysis of the influence of gas inlet delay time on interface weld ratio

Summary

The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

(1) The gas pressure loading curve based on back pressure can effectively suppress surface groove defects in four-layer structures but significantly influences their interface weld ratio. Dispersed spot-like unwelded defects are primarily distributed between the inner and outer skins that come into contact first. The triangular stiffener areas contact and form during the final stage. If the back pressure has been vented to a low level by this time, the interface weld ratio in these areas remains good.
(2) When back pressure time exceeds 40 min , gas from the outer layer cannot be vented in a timely manner and becomes trapped within closed voids. These voids appear approximately circular and are uniformly dispersed at the diffusion interface, ultimately forming unwelded defects.
(3) When back pressure time exceeds 40 min , as back pressure time increases, the amount of residual gas trapped within the voids increases, leading to a decrease in the interface weld ratio.
(4) When back pressure time exceeds 40 min , under the same back pressure time condition, an increase in either the back pressure value or the gas inlet delay time reduces the deformation rate of the inner skin and delays the contact time between the inner and outer skins (equivalent to a reduction in effective back pressure time), thereby improving the interface weld ratio.
(5) When the back pressure value is maintained between 8 and 13 bar, the back pressure time between 20 and 30 min , and the gas inlet delay time between 5 and 10 min , the fabricated four-layer structure components are free of surface grooves and exhibit an interface weld ratio of 95% or higher.

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The original version of this paper is available on https://www.scientific.net/MSF.1185.171.pdf

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Keywords Index

A

Aa6082 Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Additive Manufacturing Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Aluminium Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Aluminum Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Aluminum Alloy Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Asymmetric Tumbling Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

B

Busbars Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

C

Carbon Fiber Reinforced Polymers Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Clinched Joints A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

Cold Joining Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Continuous Compound Casting Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Current Collectors Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

D

Dissimilar Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

E

Electrical Contacting Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

End Of Life (eol) Vehicles Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Extrusion Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

F

Factory Joint Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Fe-simulation Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Fem Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine; Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Formability Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Four-layer Sandwich Structure Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Friction Stir Welding Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Fse Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Fsw Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

G

Gas Pressure Loading Curve Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

H

Heat Superimposition Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

Hybrid Materials Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

I

Interface Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Interface Weld Ratio Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Interfacial Bonding Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

J

Joining Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints; A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

Joining By Forming Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

L

L-pbf Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Lightweight Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

Lightweight Composites Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Low Cost Tooling Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

M

Magnesium Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Magnesium Sheets Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Mechanical Properties Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Micro Friction Stir Spot Welding Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Microstructure Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Movable Die Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Multi-material Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

N

Non-rotationally Symmetric Joints A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

P

Polyphenylene Sulfide Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Printed Circuit Boards Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

R

Recycling Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Refill Friction Stir Spot Welding Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Rivet Design Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Roll Bonding Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Rolling Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Rotary Friction Welding Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

S

Self-piercing Riveting Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Self-piercing Riveting (spr) Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

Sheet Metal Joining Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Silicon Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Solid State Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Solid-state Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Superplastic Forming/diffusion Bonding (spf/db) Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Surface Preparation Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Sustainable Manufacturing Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

T

Tailored Forming Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Titanium Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Tool Geometry Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Tumbling Selfpiercing Riveting (t-spr) Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

U

Ultrasonic Vibration Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

V

Versatile Joining Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

Versatility Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Void Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

W

Welding Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

X

Xlpe Insulation Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Authors Index

A

Antonini, L. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

B

Bailly, D. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Barnett, P. Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Behrens, B. Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints; Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Bobbert, M. Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Bocchi, S. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Brunotte, K. Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

Buffa, G. Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

Buonadonna, P. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

C

Carta, M. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Clark, A. Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

D

D'Urso, G. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Demirezen, A. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Devulapally, D. A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

Donati, L. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

F

Fratini, L. Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

G

Gao, J. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Giardini, C. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

Goeldner, M. Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

H

Hoyer, J. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

K

Kaimann, P. Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Kammerloher, S. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Klusemann, B. Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

L

Lechner, M. Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

Lei, H. Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Li, X. Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Lian, J. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Liu, S. Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

Liu, Z. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Loi, G. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Lv, Y. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

M

Malaske, L. Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

Mamros, E. Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Manolakos, D. A Hybrid Joining Strategy for Al-Cu Bimetallic Sheets Produced by Friction Stir Welding

Mathew, A. Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Mayer, J. Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Mehtedi, M. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Meschut, G. Numerical Analysis of the Influence of a Movable Die on Joint Formation in Versatile Self-Piercing Riveting

Mohnfeld, N. Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Mu, Y. Study on the Influence of Gas Pressure Loading Curve on Interface Weld Ratio in Superplastic Forming/Diffusion Bonding of Four-Layer Sandwich Structures

N

Negozio, M. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding

P

Papantoniou, I. A Hybrid Joining Strategy for Al-Cu Bimetallic Sheets Produced by Friction Stir Welding

Peddinghaus, J. Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints

Pelaccia, R. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Piwek, A. Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints; Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Puleo, R. Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

R

Ragai, I. Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

Reggiani, B. Influence of Rolling Temperature on the Mechanical and Formability Properties of AA1050/AZ31/AA1050 Roll-Bonded Sheets

Ringel, A. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

Russo, S. Tool Geometry Effect in Friction Stir Extrusion of Seamless Tubes from Aluminium Chips

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Sarris, J. Influence of Tumbling-Induced Geometric Anisotropy on the Mechanical Performance of Self-Piercing Rivet Joints

Schwichtenberg, N. Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Sigl, M. Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Strano, M. Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

Suhuddin, U. Effect of Surface Conditions and Tool to Interface Spacing in Refill Friction Stir Spot Welding of Aluminum to Titanium

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Tesch, M. Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Tröster, T. A Numerical Study on the Mutual Influence of Joint Orientation and Component Geometry in Non-Rotationally Symmetric Clinched Joints

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Uhe, J. Potentials of Heat Superimposed Rotary Friction Welding for Steel-Aluminium Joints; Experimental and Numerical Study of Steel-Aluminium Rotary Friction Welding with Heat Flux via Subroutine

Unni, D. Development of a Novel Rivet-Based Joining-by-Forming Process Using Flat Dies

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Volk, W. Influences of the Rolling Parameters on Multi-Material Copper-Aluminum Composites via Compound Casting

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Wang, C. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Wu, Z. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

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Yang, W. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Yin, L. Enhancing the Bonding of Insulated Homogeneous Interfaces in Factory Joints via Ultrasonic Vibration

Young, B. Effect of Fiber Length in Friction Stir Welding of Discontinuous Composites

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Zaeh, M. Micro Friction Stir Spot Welding for Copper Current Collector Applications - An Experimental Study

Zambelli, M. Experimental and Numerical Evaluation of L-PBF Printed Tool for Friction Stir Welding