In order to evaluate the benefits of controlled reduction of the billet temperature, several simulations will be carried out by changing the temperature and speed parameters. After analysing the best way to show the possibilities of the methodology, 3 sets of conditions of billet temperature, extrusion velocity and billet length have been defined. Table 1 shows the conditions chosen for the different trials.

<figure><img src="image_760.png" /><figcaption>Fig. 3. Image of the die used throughout the experiment.</figcaption></figure>

The first trial uses the conditions under which this die used to extrude in Alumex PLC. There, usually all dies extrude with a billet temperature in the range 460-470°C.

The second trial uses a billet temperature 20°C lower and doubles the extrusion speed. These are the conditions that have been optimised with the help of FE simulation to significantly increase the speed while keeping the outlet temperature and maximum press force at admissible values. The condition set for determining the temperature was not to exceed 550°C at the exit of the profile and not to approach the maximum press force (23MN) during the process.

The conditions of the third trial are chosen to show that, in the event of a possible excess of force in the press due to the temperature reduction, it is possible to reduce the billet length to reduce the maximum force required.

<table><caption>Table 1. Conditions chosen for the different research trials.</caption><thead><tr><td>Trial number</td><td>Billet temperature [°C]</td><td>Ram / extrusion velocity [mm/s]</td><td>Billet length [mm]</td></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td></tr></tbody></table>

## Numerical modelling results

Definition of the FE model. To verify the model predictions, they are compared with numerical simulation predictions employing the software Qform UK. It is a special-purpose software for the extrusion simulation based on the Lagrange–Euler approach. This software provides material flow analysis coupled with the mechanical problem in the tooling set, considering shape changes produced by the die deformation on the material flow through the die, which helps to ensure high accuracy of the numerical results [12].

All meshes are created with tetrahedral elements. A highly adaptive meshing is applied in order to optimise mesh quality and minimise the resources required to achieve accurate calculation
results. The size of the elements in each zone is determined according to the extent of local deformation undergone during the process.

The billet material is EN AW-6063-O aluminium alloy and the die material is AISI H-13 steel. All the properties of these materials are temperature-dependent functions.

With regard to 6063 aluminium alloy, density, specific heat, thermal expansion, thermal conductivity, Young's modulus and Poisson's ratio are defined by employing a linear temperature-dependent function. The H-S model is used to represent the flow stress, and its dependence on strain is also considered. The H-S model function is derived by regressing the experimental data obtained from hot torsion tests [8]. The model employed by Qform UK is based on 6 coefficients (Eq. 1).

$$ \sigma = A \cdot e^{m_1 T} \cdot \varepsilon^{m_2} \cdot e^{\frac{m_4}{\varepsilon}} \cdot (1 + \varepsilon)^{m_5 T} \cdot \varepsilon^{m_3} \quad (1) $$

The AISI H-13 steel in the die is considered to be an elastic-plastic continuum subjected to small deformations. Steel density, yield stress and Young module are temperature dependent functions. Poisson's ratio for steel is 0.3 for this alloy and is independent of temperature. In the contact areas between aluminium and steel the heat exchange coefficient is 30000 W/(m²·K).

<table><caption>Table 2. Temperatures of tooling.</caption><thead><tr><th></th><th>Temp [°C]</th><th></th><th>Temp [°C]</th></tr></thead><tbody><tr><td>Container</td><td>440</td><td>Bolster</td><td>280</td></tr><tr><td>Ram</td><td>450</td><td>Die</td><td>450</td></tr><tr><td>Die ring</td><td>450</td><td>Backer</td><td>450</td></tr></tbody></table>

Qform UK models the friction between the aluminium and the die surface using the friction model proposed by Levanov [13], where $\sigma_n$ is the normal contact pressure and m is the friction factor (Eq. 2):

$$ f_{\tau} = m \cdot \bar{\sigma} / \sqrt{3} \left[ 1 - \exp \left( -1.25 \cdot \frac{\sigma_n}{\bar{\sigma}} \right) \right] \quad (2) $$

In addition to the boundary conditions to be varied during the experiment, listed in Table 1, Table 2 shows the remaining main boundary conditions of the extrusion process used in the Qform UK FEM simulation.

<table><caption>Table 3. Numerical results for the different virtual trials.</caption><thead><tr><th>Trial number</th><th>Billet temp. [°C]</th><th>Ram velocity [mm/s]</th><th>Billet length [mm]</th><th>Profile temp. [°C]</th><th>Max pressure [bar]</th><th>Max velocity deviation [%]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>516</td><td>188.6</td><td>11.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>548</td><td>206.3</td><td>11</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>546</td><td>183.4</td><td>10.5</td></tr></tbody></table>

Model results. From all the results available in the simulation, those that are most representative for the research have been chosen:

- Temperature at process outlet (°C). It allows to control the temperature of the extruded profile at the exit of the die.

- Absolute value of the velocity (mm/s). It allows to show the velocity reached by the process in steady state.

- Ram extrusion load (MN) and maximal pressure (bar). It allows to control the maximum force to be exerted by the ram of the press and the maximal pressure during the extrusion.
results. The size of the elements in each zone is determined according to the extent of local deformation undergone during the process.

The billet material is EN AW-6063-O aluminium alloy and the die material is AISI H-13 steel. All the properties of these materials are temperature-dependent functions.

With regard to 6063 aluminium alloy, density, specific heat, thermal expansion, thermal conductivity, Young's modulus and Poisson's ratio are defined by employing a linear temperature-dependent function. The H-S model is used to represent the flow stress, and its dependence on strain is also considered. The H-S model function is derived by regressing the experimental data obtained from hot torsion tests [8]. The model employed by Qform UK is based on 6 coefficients (Eq. 1).

$$ \sigma = A \cdot e^{m_1 T} \cdot \varepsilon^{m_2} \cdot e^{\frac{m_4}{\varepsilon}} \cdot (1 + \varepsilon)^{m_5 T} \cdot \varepsilon^{m_3} \quad (1) $$

The AISI H-13 steel in the die is considered to be an elastic-plastic continuum subjected to small deformations. Steel density, yield stress and Young module are temperature dependent functions. Poisson's ratio for steel is 0.3 for this alloy and is independent of temperature. In the contact areas between aluminium and steel the heat exchange coefficient is 30000 W/(m²·K).

<table><caption>Table 2. Temperatures of tooling.</caption><thead><tr><th></th><th>Temp [°C]</th><th></th><th>Temp [°C]</th></tr></thead><tbody><tr><td>Container</td><td>440</td><td>Bolster</td><td>280</td></tr><tr><td>Ram</td><td>450</td><td>Die</td><td>450</td></tr><tr><td>Die ring</td><td>450</td><td>Backer</td><td>450</td></tr></tbody></table>

Qform UK models the friction between the aluminium and the die surface using the friction model proposed by Levanov [13], where $\sigma_n$ is the normal contact pressure and m is the friction factor (Eq. 2):

$$ f_{\tau} = m \cdot \bar{\sigma} / \sqrt{3} \left[ 1 - \exp(-1.25 \cdot \sigma_{n} / \bar{\sigma}) \right] \quad (2) $$

In addition to the boundary conditions to be varied during the experiment, listed in Table 1, Table 2 shows the remaining main boundary conditions of the extrusion process used in the Qform UK FEM simulation.

<table><caption>Table 3. Numerical results for the different virtual trials.</caption><thead><tr><th>Trial number</th><th>Billet temp. [°C]</th><th>Ram velocity [mm/s]</th><th>Billet length [mm]</th><th>Profile temp. [°C]</th><th>Max pressure [bar]</th><th>Max velocity deviation [%]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>516</td><td>188.6</td><td>11.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>548</td><td>206.3</td><td>11</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>546</td><td>183.4</td><td>10.5</td></tr></tbody></table>

Model results. From all the results available in the simulation, those that are most representative for the research have been chosen:

- Temperature at process outlet (°C). It allows to control the temperature of the extruded profile at the exit of the die.

- Absolute value of the velocity (mm/s). It allows to show the velocity reached by the process in steady state.

- Ram extrusion load (MN) and maximal pressure (bar). It allows to control the maximum force to be exerted by the ram of the press and the maximal pressure during the extrusion.
- Velocity deviation (%). It allows to control that the flow during the process is balanced comparing the velocity deviation between the fastest and slowest point in the extruded profile.

<figure><img src="image_753.png" /><figcaption>Fig. 4. Exit temperature and absolute velocity results for trial 1.</figcaption></figure>

<figure><img src="image_754.png" /><figcaption>Fig. 5. Velocity deviation and ram load results for trial 1.</figcaption></figure>

<figure><img src="image_755.png" /><figcaption>Fig. 6. Exit temperature and absolute velocity results for trial 2.</figcaption></figure>
- Velocity deviation (%). It allows to control that the flow during the process is balanced comparing the velocity deviation between the fastest and slowest point in the extruded profile.

<figure><img src="image_761.png" /><figcaption>Fig. 4. Exit temperature and absolute velocity results for trial 1.</figcaption></figure>

<figure><img src="image_762.png" /><figcaption>Fig. 5. Velocity deviation and ram load results for trial 1.</figcaption></figure>

<figure><img src="image_763.png" /><figcaption>Fig. 6. Exit temperature and absolute velocity results for trial 2.</figcaption></figure>
<figure><img src="image_756.png" /><figcaption>Fig. 7. Velocity deviation and ram load results for trial 2.</figcaption></figure>

<figure><img src="image_757.png" /><figcaption>Fig. 8. Exit temperature and velocity deviation results for trial 3.</figcaption></figure>

<figure><img src="image_758.png" /><figcaption>Fig. 9. Velocity deviation and ram load results for trial 3.</figcaption></figure>

For each of the trials performed, the results obtained in the FE analyses are shown in Table 3. Fig. 4 to Fig. 9 show the exit temperature, absolute profile velocity, velocity deviation and ram load results for each of the trials. The rest of the graphs of the results can be consulted in the dataset associated with the research [14].
<figure><img src="image_764.png" /><figcaption>Fig. 7. Velocity deviation and ram load results for trial 2.</figcaption></figure>

<figure><img src="image_765.png" /><figcaption>Fig. 8. Exit temperature and velocity deviation results for trial 3.</figcaption></figure>

<figure><img src="image_766.png" /><figcaption>Fig. 9. Velocity deviation and ram load results for trial 3.</figcaption></figure>

For each of the trials performed, the results obtained in the FE analyses are shown in Table 3. Fig. 4 to Fig. 9 show the exit temperature, absolute profile velocity, velocity deviation and ram load results for each of the trials. The rest of the graphs of the results can be consulted in the dataset associated with the research [14].
## Trials results

The three trials have been carried out under exactly the same conditions as those defined in the simulations described above. It should be noted that in all trials the extrusion result was correct:

- The flow of the aluminium was sufficiently balanced.

- The surface finishing of the profiles is correct.

- The profiles conform to all dimensional tolerances.

Fig. 10 shows the images of the test specimens resulting from trials 1 and 2. The remaining images of the trials and the videos of the process can be consulted in the dataset corresponding to this paper [14]. Table 4 shows the actual results obtained in each of the trials in the press.

<table><caption>Table 4. Actual results for the different research trials.</caption><thead><tr><th>Trial number</th><th>Billet temp. [°C]</th><th>Ram velocity [mm/s]</th><th>Billet length [mm]</th><th>Profile temp. [°C]</th><th>Max pressure [bar]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>523</td><td>200.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>565</td><td>208.8</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>564</td><td>195.3</td></tr></tbody></table>

<figure><img src="image_759.png" /><figcaption>Fig. 10. Test specimens resulting from trial 1 (left) and trial 2 (right).</figcaption></figure>

## Discussion

Table 5 shows the comparison of the experimental results versus the FE simulation results. An analysis of the values confronted in this table shows that, in general terms, they are very similar. The profile temperature values are very slightly higher in the real trials (2-3%). The pressure values show a somewhat larger difference (up to 6.5% in the case of trials 1 and 3), being always higher in the real trials.

<table><caption>Table 5. Comparison of the experimental results versus the FE simulation results.</caption><thead><tr><th>Trial</th><th>Billet temp.[°C]</th><th>Ram vel. [mm/s]</th><th>Billet [mm]</th><th>Profile temp. FE Simul. [°C]</th><th>Profile temp. Press [°C]</th><th>Max pressure FE Simul.[bar]</th><th>Max pressure Press [bar]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>516</td><td>523</td><td>188.6</td><td>200.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>548</td><td>565</td><td>206.3</td><td>208.8</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>546</td><td>564</td><td>183.4</td><td>195.3</td></tr></tbody></table>

Therefore, it is possible to state that the predictions of the FE simulation are very close to reality and are reliable for making decisions to reduce the billet temperature in order to increase the velocity of the process. This is the most important point as the proposed methodology is based on
## Trials results

The three trials have been carried out under exactly the same conditions as those defined in the simulations described above. It should be noted that in all trials the extrusion result was correct:

- The flow of the aluminium was sufficiently balanced.

- The surface finishing of the profiles is correct.

- The profiles conform to all dimensional tolerances.

Fig. 10 shows the images of the test specimens resulting from trials 1 and 2. The remaining images of the trials and the videos of the process can be consulted in the dataset corresponding to this paper [14]. Table 4 shows the actual results obtained in each of the trials in the press.

<table><caption>Table 4. Actual results for the different research trials.</caption><thead><tr><th>Trial number</th><th>Billet temp. [°C]</th><th>Ram velocity [mm/s]</th><th>Billet length [mm]</th><th>Profile temp. [°C]</th><th>Max pressure [bar]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>523</td><td>200.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>565</td><td>208.8</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>564</td><td>195.3</td></tr></tbody></table>

<figure><img src="image_767.png" /><figcaption>Fig. 10. Test specimens resulting from trial 1 (left) and trial 2 (right).</figcaption></figure>

## Discussion

Table 5 shows the comparison of the experimental results versus the FE simulation results. An analysis of the values confronted in this table shows that, in general terms, they are very similar. The profile temperature values are very slightly higher in the real trials (2-3%). The pressure values show a somewhat larger difference (up to 6.5% in the case of trials 1 and 3), being always higher in the real trials.

<table><caption>Table 5. Comparison of the experimental results versus the FE simulation results.</caption><thead><tr><th>Trial</th><th>Billet temp.[°C]</th><th>Ram vel. [mm/s]</th><th>Billet [mm]</th><th>Profile temp. FE Simul. [°C]</th><th>Profile temp. Press [°C]</th><th>Max pressure FE Simul.[bar]</th><th>Max pressure Press [bar]</th></tr></thead><tbody><tr><td>1</td><td>467</td><td>3</td><td>700</td><td>516</td><td>523</td><td>188.6</td><td>200.5</td></tr><tr><td>2</td><td>447</td><td>6</td><td>700</td><td>548</td><td>565</td><td>206.3</td><td>208.8</td></tr><tr><td>3</td><td>442</td><td>6</td><td>520</td><td>546</td><td>564</td><td>183.4</td><td>195.3</td></tr></tbody></table>

Therefore, it is possible to state that the predictions of the FE simulation are very close to reality and are reliable for making decisions to reduce the billet temperature in order to increase the velocity of the process. This is the most important point as the proposed methodology is based on
the use of FE simulation prediction to directly choose the most appropriate conditions to maximise press productivity without risk or the need for iterative testing at different temperatures.

On the other hand, comparing the results of the first two trials, it should be noted that, as expected, with a reduction of 20°C in the billet temperature it has been possible to double the process speed with an increase in profile temperature of only 42°C (maintained at 565°C, within an admissible range, as shown by the good surface finish of the specimens).

Comparing the results of trials 2 and 3, it can also be seen, as expected, that by reducing the billet length it is possible to reduce the maximum pressure/force exerted by the press.

From the FE simulation results it can be seen that the maximum predicted speed deviations are around 10.5-11.5%. The validity of this maximum deviation is highly dependent on the geometry and tolerances of the profile. As it has been verified later on the press, they are totally admissible and do not pose any problem in obtaining the extruded profile. It is observed that the velocity deviations are distributed exactly the same for the different process conditions analysed, which indicates that the creep behaviour of the extruded profile in the press is very similar for all cases.

## Conclusions and future research

The main conclusion drawn is that the proposed methodology to optimise the extrusion parameters by means of FE simulation is very valuable from the point of view that it allows deciding the optimum extrusion conditions to maximise the productivity of the press without any risk for the tools and without the need to carry out iterative trials on the press, with the cost that this involves.

Furthermore, the numerical and experimental analysis conducted in this study shows that the strategy of reducing the billet temperature is a good solution for increasing the productivity of some dies. Moreover, it is a strategy that does not require investment in press equipment, compared to other strategies such as nitrogen cooling.

It shows that it is important to discard the widespread strategy of using standardised and equal conditions for the extrusion of most dies because potential process productivity is being wasted by doing so.

Finally, the question arises as to whether it is cost-effective to use simulation to try to increase the productivity of the extrusion press. From the authors' point of view, for companies that manufacture their own dies and/or purchase dies that have been designed using simulation, the cost of this type of study for each die is very low and the potential benefit is large. In these cases, moreover, it is a change that can be introduced in the very short term. And since these are process changes that do not require any investment in press equipment, any improvement in productivity is always positive.

For companies that are totally unfamiliar with FE simulation, this strategy is more difficult to take on board in the short term. It requires a medium-term strategy to introduce FE simulation into the corporate culture.

This billet temperature reduction methodology must always be accompanied by simulation to control the final profile temperature and maximum press pressures. In any case, it would be possible to avoid its use for dies very similar to others already studied.

For future research, it is intended to further develop this methodology by analysing cases of more complex solid dies and hollow dies. Alumex PLC is a perfect candidate to use the methodology and further develop it because it has fully integrated FE simulation in the development and design of its dies.

## References

[1] A. F. M. Arif, A. K. Sheikh, & S. Z. Qamar, A study of die failure mechanisms in aluminum extrusion. *Journal of Materials Processing Technology*, 134 (2003) 318-328.
https://doi.org/10.1016/S0924-0136(02)01116-0
the use of FE simulation prediction to directly choose the most appropriate conditions to maximise press productivity without risk or the need for iterative testing at different temperatures.

On the other hand, comparing the results of the first two trials, it should be noted that, as expected, with a reduction of 20°C in the billet temperature it has been possible to double the process speed with an increase in profile temperature of only 42°C (maintained at 565°C, within an admissible range, as shown by the good surface finish of the specimens).

Comparing the results of trials 2 and 3, it can also be seen, as expected, that by reducing the billet length it is possible to reduce the maximum pressure/force exerted by the press.

From the FE simulation results it can be seen that the maximum predicted speed deviations are around 10.5-11.5%. The validity of this maximum deviation is highly dependent on the geometry and tolerances of the profile. As it has been verified later on the press, they are totally admissible and do not pose any problem in obtaining the extruded profile. It is observed that the velocity deviations are distributed exactly the same for the different process conditions analysed, which indicates that the creep behaviour of the extruded profile in the press is very similar for all cases.

## Conclusions and future research

The main conclusion drawn is that the proposed methodology to optimise the extrusion parameters by means of FE simulation is very valuable from the point of view that it allows deciding the optimum extrusion conditions to maximise the productivity of the press without any risk for the tools and without the need to carry out iterative trials on the press, with the cost that this involves.

Furthermore, the numerical and experimental analysis conducted in this study shows that the strategy of reducing the billet temperature is a good solution for increasing the productivity of some dies. Moreover, it is a strategy that does not require investment in press equipment, compared to other strategies such as nitrogen cooling.

It shows that it is important to discard the widespread strategy of using standardised and equal conditions for the extrusion of most dies because potential process productivity is being wasted by doing so.

Finally, the question arises as to whether it is cost-effective to use simulation to try to increase the productivity of the extrusion press. From the authors' point of view, for companies that manufacture their own dies and/or purchase dies that have been designed using simulation, the cost of this type of study for each die is very low and the potential benefit is large. In these cases, moreover, it is a change that can be introduced in the very short term. And since these are process changes that do not require any investment in press equipment, any improvement in productivity is always positive.

For companies that are totally unfamiliar with FE simulation, this strategy is more difficult to take on board in the short term. It requires a medium-term strategy to introduce FE simulation into the corporate culture.

This billet temperature reduction methodology must always be accompanied by simulation to control the final profile temperature and maximum press pressures. In any case, it would be possible to avoid its use for dies very similar to others already studied.

For future research, it is intended to further develop this methodology by analysing cases of more complex solid dies and hollow dies. Alumex PLC is a perfect candidate to use the methodology and further develop it because it has fully integrated FE simulation in the development and design of its dies.

## References

[1] A. F. M. Arif, A. K. Sheikh, & S. Z. Qamar, A study of die failure mechanisms in aluminum extrusion. *Journal of Materials Processing Technology*, 134 (2003) 318-328.
https://doi.org/10.1016/S0924-0136(02)01116-0
[2] E. Giarmas & D. Tzetzis, Optimization of die design for extrusion of 6xxx series aluminum alloys through finite element analysis: a critical review. *International Journal of Advanced Manufacturing Technology*, 119 (2022) 5529–5551. https://doi.org/10.1007/s00170-022-08694-3

[3] S. S. Akhtar & A. F. M. Arif, Fatigue failure of extrusion dies: Effect of process parameters and design features on die life. *Journal of Failure Analysis and Prevention*, 10 (2010). https://doi.org/10.1007/s11668-009-9304-4

[4] S. Z. Qamar, A. K. Sheikh, A. F. M. Arif, M. Younas, & T. Pervez, Monte Carlo simulation of extrusion die life. *Journal of Materials Processing Technology*, (2008). https://doi.org/10.1016/j.jmatprotec.2007.08.062

[5] S. Z. Qamar, A. F. M. Arif, & A. K. Sheikh, Analysis of product defects in a typical aluminum extrusion facility. *Materials and Manufacturing Processes*, 19 (2004). https://doi.org/10.1081/AMP-120038650

[6] P. K. Saha, Thermodynamics and tribology in aluminum extrusion. *Wear*, 218 (1998). https://doi.org/10.1016/S0043-1648(98)00210-5

[7] R. Pelaccia, B. Reggiani, M. Negozio, & L. Donati, Liquid nitrogen in the industrial practice of hot aluminium extrusion: experimental and numerical investigation. *International Journal of Advanced Manufacturing Technology*, 119 (2022). https://doi.org/10.1007/s00170-021-08422-3

[8] A. Gamberoni, L. Donati, B. Reggiani, M. Haase, L. Tomesani, & A. E. Tekkaya, Industrial Benchmark 2015: Process Monitoring and Analysis of Hollow EN AW-6063 Extruded Profile. *Materials Today: Proceedings*, 2 (2015) 4714–4725. https://doi.org/10.1016/j.matpr.2015.10.004

[9] H. H. Jo, S. K. Lee, C. S. Jung, & B. M. Kim, A non-steady state FE analysis of Al tubes hot extrusion by a porthole die. *Journal of Materials Processing Technology*, 173 (2006). https://doi.org/10.1016/j.jmatprotec.2005.03.039

[10] A. Medvedev, A. Bevacqua, A. Molotnikov, R. Axe, & R. Lapovok, Innovative aluminium extrusion: Increased productivity through simulation. *Procedia Manuf.* (2020). https://doi.org/10.1016/j.promfg.2020.08.085

[11] M. M. Marín, A. M. Camacho, & J. A. Pérez, Influence of the temperature on AA6061 aluminum alloy in a hot extrusion process. *Procedia Manufacturing*, 13 (2017) 327–334. https://doi.org/10.1016/j.promfg.2017.09.084

[12] N. Biba, S. Stebunov, & A. Lishny, The model for coupled simulation of thin profile extrusion. *Key Engineering Materials*, 504–506 (2012) 505–510. https://doi.org/10.4028/www.scientific.net/KEM.504-506.505

[13] A. N. Levanov, Improvement of metal forming processes by means of useful effects of plastic friction. *Journal of Materials Processing Technology*, 72 (1997) 314–316. https://doi.org/10.1016/S0924-0136(97)00191-X

[14] J. Llorca-Schenk, L. Gayashan, & P. Hewage, Dataset for increasing of productivity of aluminum hot extrusion press. *Mendeley Data*, (2024). https://doi.org/10.17632/chr2mnfh3j.1
[2] E. Giarmas & D. Tzetzis, Optimization of die design for extrusion of 6xxx series aluminum alloys through finite element analysis: a critical review. *International Journal of Advanced Manufacturing Technology*, 119 (2022) 5529–5551. https://doi.org/10.1007/s00170-022-08694-3

[3] S. S. Akhtar & A. F. M. Arif, Fatigue failure of extrusion dies: Effect of process parameters and design features on die life. *Journal of Failure Analysis and Prevention*, 10 (2010). https://doi.org/10.1007/s11668-009-9304-4

[4] S. Z. Qamar, A. K. Sheikh, A. F. M. Arif, M. Younas, & T. Pervez, Monte Carlo simulation of extrusion die life. *Journal of Materials Processing Technology*, (2008). https://doi.org/10.1016/j.jmatprotec.2007.08.062

[5] S. Z. Qamar, A. F. M. Arif, & A. K. Sheikh, Analysis of product defects in a typical aluminum extrusion facility. *Materials and Manufacturing Processes*, 19 (2004). https://doi.org/10.1081/AMP-120038650

[6] P. K. Saha, Thermodynamics and tribology in aluminum extrusion. *Wear*, 218 (1998). https://doi.org/10.1016/S0043-1648(98)00210-5

[7] R. Pelaccia, B. Reggiani, M. Negozio, & L. Donati, Liquid nitrogen in the industrial practice of hot aluminium extrusion: experimental and numerical investigation. *International Journal of Advanced Manufacturing Technology*, 119 (2022). https://doi.org/10.1007/s00170-021-08422-3

[8] A. Gamberoni, L. Donati, B. Reggiani, M. Haase, L. Tomesani, & A. E. Tekkaya, Industrial Benchmark 2015: Process Monitoring and Analysis of Hollow EN AW-6063 Extruded Profile. *Materials Today: Proceedings*, 2 (2015) 4714–4725. https://doi.org/10.1016/j.matpr.2015.10.004

[9] H. H. Jo, S. K. Lee, C. S. Jung, & B. M. Kim, A non-steady state FE analysis of Al tubes hot extrusion by a porthole die. *Journal of Materials Processing Technology*, 173 (2006). https://doi.org/10.1016/j.jmatprotec.2005.03.039

[10] A. Medvedev, A. Bevacqua, A. Molotnikov, R. Axe, & R. Lapovok, Innovative aluminium extrusion: Increased productivity through simulation. *Procedia Manuf.* (2020). https://doi.org/10.1016/j.promfg.2020.08.085

[11] M. M. Marín, A. M. Camacho, & J. A. Pérez, Influence of the temperature on AA6061 aluminum alloy in a hot extrusion process. *Procedia Manufacturing*, 13 (2017) 327–334. https://doi.org/10.1016/j.promfg.2017.09.084

[12] N. Biba, S. Stebunov, & A. Lishny, The model for coupled simulation of thin profile extrusion. *Key Engineering Materials*, 504–506 (2012) 505–510. https://doi.org/10.4028/www.scientific.net/KEM.504-506.505

[13] A. N. Levanov, Improvement of metal forming processes by means of useful effects of plastic friction. *Journal of Materials Processing Technology*, 72 (1997) 314–316. https://doi.org/10.1016/S0924-0136(97)00191-X

[14] J. Llorca-Schenk, L. Gayashan, & P. Hewage, Dataset for increasing of productivity of aluminum hot extrusion press. *Mendeley Data*, (2024). https://doi.org/10.17632/chr2mnfh3j.1
Forging and rolling

<figure><img src="image_760.png" /></figure>
Forging and rolling

<figure><img src="image_768.png" /></figure>
# Influence of enlarged joining zone interfaces on the bond properties of tailored formed hybrid components made of 20MnCr5 steel and EN AW-6082 aluminium

PIWEK Armin¹,ᵃ*, PEDDINGHAUS Julius¹,ᵇ, UHE Johanna¹,ᶜ and BRUNOTTE Kai¹,ᵈ

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

**Keywords:** Friction Welding, Hot Forging, Tailored Forming, Joining Zone, Spray Cooling

**Abstract.** Hybrid material compounds offer an extension of the technological application range of monolithic components by combining positive material-specific properties. In the case of steel and aluminium, a load-adapted component with high strength areas and a reduced weight can be created. Tailored Forming enables the joining zone created by a pre-joining process to be modified and enhanced by a subsequent forming step. Derived from previous studies, an enlarged joining zone interface through spherical joining zone curvature and an equalisation of yield stresses through an inhomogeneous induction heating with partial cooling are necessary to achieve a defect free bond with high strength and ductility. In order to further enlarge the joining zone interface and hence to increase the surface ratio of juvenile welding spots without brittle intermetallic compounds, different local plastic strains are induced. Additionally, an alternative spray cooling concept is used to evaluate the effect of steeper temperature gradients on the bond quality. Rotary friction welded specimens made of 20MnCr5 steel and EN AW-6082 aluminium are cup backward extruded with different extrusion ratios using punch diameters of 22 mm and 30 mm. Metallographic images, SEM analysis and hardness tests of cross-sections are used to evaluate the bond quality with regard to the joining zone formation, occurring defects and the resulting intermetallic compound. With cooling, higher yield stresses could be set in the aluminium, which counteract material failure even with larger punch diameters due to a higher deformability. However, the surface enlargement of the joining zone is reduced. Despite the higher surface enlargement in uncooled specimens, insufficient bonds were achieved due to existing cracks in the aluminium in or near the joining zone interface, as well as significant thicker intermetallic compounds.

**Introduction**

In a world characterised by constant progress and innovation, lightweight construction is becoming progressively important to optimise technical high-performance components. This paradigm shift in industrial applications not only promises a weight reduction, but also a load-adapted design to overcome the limitations of monolithic components. Hybrid material concepts enable structures to be designed to conserve resources and also opens up new horizons for sustainable mobility, increased efficiency as well as a variety of applications in a wide range of industries. The combination of different material properties is already a well-established practice in sheet metal forming, commonly referred to as tailored blanks, and is extensively applied in the construction of automobile bodies. Joined sheets created from steel with varying qualities, thicknesses, strength or surface coatings are shaped in a final step to create specific elements e. g. b-pillars with high strength and also sufficient elongation capacity to ensure safety in the event of a crash [1].

<figure><img src="image_761.png" /></figure>

<footnote>ᵃpiwek@ifum.uni-hannover.de, ᵇpeddinghaus@ifum.uni-hannover.de,</footnote>

<footnote>ᶜuhe@ifum.uni-hannover.de, ᵈbrunotte@ifum.uni-hannover.de</footnote>
# Influence of enlarged joining zone interfaces on the bond properties of tailored formed hybrid components made of 20MnCr5 steel and EN AW-6082 aluminium

PIWEK Armin¹,ᵃ*, PEDDINGHAUS Julius¹,ᵇ, UHE Johanna¹,ᶜ
and BRUNOTTE Kai¹,ᵈ

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

**Keywords:** Friction Welding, Hot Forging, Tailored Forming, Joining Zone, Spray Cooling

**Abstract.** Hybrid material compounds offer an extension of the technological application range of monolithic components by combining positive material-specific properties. In the case of steel and aluminium, a load-adapted component with high strength areas and a reduced weight can be created. Tailored Forming enables the joining zone created by a pre-joining process to be modified and enhanced by a subsequent forming step. Derived from previous studies, an enlarged joining zone interface through spherical joining zone curvature and an equalisation of yield stresses through an inhomogeneous induction heating with partial cooling are necessary to achieve a defect free bond with high strength and ductility. In order to further enlarge the joining zone interface and hence to increase the surface ratio of juvenile welding spots without brittle intermetallic compounds, different local plastic strains are induced. Additionally, an alternative spray cooling concept is used to evaluate the effect of steeper temperature gradients on the bond quality. Rotary friction welded specimens made of 20MnCr5 steel and EN AW-6082 aluminium are cup backward extruded with different extrusion ratios using punch diameters of 22 mm and 30 mm. Metallographic images, SEM analysis and hardness tests of cross-sections are used to evaluate the bond quality with regard to the joining zone formation, occurring defects and the resulting intermetallic compound. With cooling, higher yield stresses could be set in the aluminium, which counteract material failure even with larger punch diameters due to a higher deformability. However, the surface enlargement of the joining zone is reduced. Despite the higher surface enlargement in uncooled specimens, insufficient bonds were achieved due to existing cracks in the aluminium in or near the joining zone interface, as well as significant thicker intermetallic compounds.

**Introduction**

In a world characterised by constant progress and innovation, lightweight construction is becoming progressively important to optimise technical high-performance components. This paradigm shift in industrial applications not only promises a weight reduction, but also a load-adapted design to overcome the limitations of monolithic components. Hybrid material concepts enable structures to be designed to conserve resources and also opens up new horizons for sustainable mobility, increased efficiency as well as a variety of applications in a wide range of industries. The combination of different material properties is already a well-established practice in sheet metal forming, commonly referred to as tailored blanks, and is extensively applied in the construction of automobile bodies. Joined sheets created from steel with varying qualities, thicknesses, strength or surface coatings are shaped in a final step to create specific elements e. g. b-pillars with high strength and also sufficient elongation capacity to ensure safety in the event of a crash [1].

<figure><img src="image_769.png" /></figure>

<footnote>ᵃpiwek@ifum.uni-hannover.de, ᵇpeddinghaus@ifum.uni-hannover.de,</footnote>

<footnote>ᶜuhe@ifum.uni-hannover.de, ᵈbrunotte@ifum.uni-hannover.de</footnote>
## State of the Art

The concept of hybrid components has already been applied to bulk metal in numerous studies in terms of compound forging, where a joint is created between two semi-finished products by forming itself [2]. Compound forging requires a certain plastic deformation and consequently an increase in surface area in order to create a sufficient bond. In the case of a bonding between steel and aluminium, oxide layers are broken up according to the Bay model so that metallic contacts and consequently covalent Al-Fe compounds can subsequently form along the interface [3]. In the studies by Groche et al. [4] and Peter et al. [5], the steel alloy C15 and the age-hardenable wrought aluminum alloy EN AW-6082 were joined by cold forming using a cup backward extrusion process. The surface enlargement of the joining zone, which has a positive effect on the bond due to oxide cracking, could be significantly influenced by the heat treatment condition and the surface roughness of the different materials. Alternatively, it is also possible to set a temperature gradient between the bonding partners, which leads to an equalisation of the different yield stresses during forming. By varying the workpiece temperatures for steel at 600 °C and aluminium at 200 °C, Wohletz and Groche [6] achieved equal or even higher bond strengths in comparison to cold compound forging. Wang et al. [7] studied the deformation mechanisms and mechanical properties in aluminium-steel bimetallic gears produced by hot compound forging. The temperature of the aluminium core was varied between 300 and 500 °C. The steel, which was formed into the external gear rim around the core, was at a temperature of 850 °C. At higher temperatures, it is more difficult to break up the oxides, as the softened aluminium induces insufficient severe plastic deformation. The best bond strengths were achieved at an aluminium temperature of 300 °C, as the oxides were broken up and pushed into the aluminium base material due to the sufficient strength of the aluminium resulting in an oxide-free interface area.

Further investigations were carried out using a combined manufacturing process for hybrid materials consisting of pre-joining and subsequent forming at elevated temperatures. In this case, the forming is not used to join the semi-finished products, but to geometrically influence the component to gain a final contour preform and in particular to thermomechanically enhance the joining zone properties. In the studies of Behrens et al. [8] or Foydl et al. [9], a combined process of extrusion and die forging was developed to enhance the mechanical properties of co-extruded profiles by embedding reinforcing steel fibres in an aluminium matrix material. However, it is also possible to create load adapted components with strictly divided areas made of different materials with complementary properties. In this regard, Domblesky et al. [10] joined hybrid semi-finished products made of copper and steel or copper and aluminium with rotary friction welding in a serial arrangement. Subsequently, the specimens were forged under hot working conditions. The hybrid parts showed a sufficient workability and are comparable to the base materials in terms of yield and tensile strength. Various studies have applied this novel Tailored Forming process chain to produce hybrid high-performance components with load-adapted areas made of steel and aluminium [11, 12]. By using pre-joined components with a material bond in a subsequent forming step, the aforementioned problem of oxides in the joining zone can be avoided and sufficiently high bond strengths can be achieved. Rotary friction welding has proven to be a suitable joining process, in which the strength-reducing oxides can be removed from the joining zone by the plastic deformation of aluminium that occurs during the process [13]. Excessive growth of the brittle intermetallic compound (IMC), which forms in the interface due to the insolubility between steel and aluminium, is not initiated by the low heat input. According to Herbst et al. [14], seam thicknesses of 0.07 µm occur in friction welded specimens made of 20MnCr5 and EN AW-6082, which are well below the critical values of 1 µm, that significantly reduce the bond strength. It is assumed that the existence of an IMC is necessary for a sufficient bond. Maximum bond strengths are achieved with an IMC thickness of 0.5 µm in several investigations [15].
## State of the Art

The concept of hybrid components has already been applied to bulk metal in numerous studies in terms of compound forging, where a joint is created between two semi-finished products by forming itself [2]. Compound forging requires a certain plastic deformation and consequently an increase in surface area in order to create a sufficient bond. In the case of a bonding between steel and aluminium, oxide layers are broken up according to the Bay model so that metallic contacts and consequently covalent Al-Fe compounds can subsequently form along the interface [3]. In the studies by Groche et al. [4] and Peter et al. [5], the steel alloy C15 and the age-hardenable wrought aluminum alloy EN AW-6082 were joined by cold forming using a cup backward extrusion process. The surface enlargement of the joining zone, which has a positive effect on the bond due to oxide cracking, could be significantly influenced by the heat treatment condition and the surface roughness of the different materials. Alternatively, it is also possible to set a temperature gradient between the bonding partners, which leads to an equalisation of the different yield stresses during forming. By varying the workpiece temperatures for steel at 600 °C and aluminium at 200 °C, Wohletz and Groche [6] achieved equal or even higher bond strengths in comparison to cold compound forging. Wang et al. [7] studied the deformation mechanisms and mechanical properties in aluminium-steel bimetallic gears produced by hot compound forging. The temperature of the aluminium core was varied between 300 and 500 °C. The steel, which was formed into the external gear rim around the core, was at a temperature of 850 °C. At higher temperatures, it is more difficult to break up the oxides, as the softened aluminium induces insufficient severe plastic deformation. The best bond strengths were achieved at an aluminium temperature of 300 °C, as the oxides were broken up and pushed into the aluminium base material due to the sufficient strength of the aluminium resulting in an oxide-free interface area.

Further investigations were carried out using a combined manufacturing process for hybrid materials consisting of pre-joining and subsequent forming at elevated temperatures. In this case, the forming is not used to join the semi-finished products, but to geometrically influence the component to gain a final contour preform and in particular to thermomechanically enhance the joining zone properties. In the studies of Behrens et al. [8] or Foydl et al. [9], a combined process of extrusion and die forging was developed to enhance the mechanical properties of co-extruded profiles by embedding reinforcing steel fibres in an aluminium matrix material. However, it is also possible to create load adapted components with strictly divided areas made of different materials with complementary properties. In this regard, Domblesky et al. [10] joined hybrid semi-finished products made of copper and steel or copper and aluminium with rotary friction welding in a serial arrangement. Subsequently, the specimens were forged under hot working conditions. The hybrid parts showed a sufficient workability and are comparable to the base materials in terms of yield and tensile strength. Various studies have applied this novel Tailored Forming process chain to produce hybrid high-performance components with load-adapted areas made of steel and aluminium [11, 12]. By using pre-joined components with a material bond in a subsequent forming step, the aforementioned problem of oxides in the joining zone can be avoided and sufficiently high bond strengths can be achieved. Rotary friction welding has proven to be a suitable joining process, in which the strength-reducing oxides can be removed from the joining zone by the plastic deformation of aluminium that occurs during the process [13]. Excessive growth of the brittle intermetallic compound (IMC), which forms in the interface due to the insolubility between steel and aluminium, is not initiated by the low heat input. According to Herbst et al. [14], seam thicknesses of 0.07 µm occur in friction welded specimens made of 20MnCr5 and EN AW-6082, which are well below the critical values of 1 µm, that significantly reduce the bond strength. It is assumed that the existence of an IMC is necessary for a sufficient bond. Maximum bond strengths are achieved with an IMC thickness of 0.5 µm in several investigations [15].
Different extrusion processes were developed and investigated with regard to the bond strength and ductility [16]. Cup backward extrusion proved to be the most suitable process for producing a hybrid hollow shaft due to the greatest influence on the joining zone geometry and consequently the mechanical properties [17]. By using an inductive heating concept with an additional immersion cooling, the temperature gradient between steel and aluminium could be significantly increased in a full forward extrusion process, which led to an improved equalisation of the yield stresses. In this case, the aluminium did not fail in the process, which is why cracks along the joining zone could be largely prevented [16]. Furthermore, in the course of this, an alternative spray cooling was theoretically designed for impact extrusion in the study by Ince et al. [18]. With this concept, greater temperature gradients can be achieved by avoiding the Leidenfrost effect.

Based on the findings presented, the combination of cup backward extrusion combined with an enhanced cooling strategy to achieve a beneficial forming temperature gradient shows high capability to utilise the full potential of the Tailored Forming process chain. This study focuses on cup backward extrusion at elevated forming temperature with partial spray cooling to study the interactions between surface enlargement, the emerging IMC and the resulting interface.

## Materials & Methods

As in previous studies [15], the material pairing of steel (soft-annealed 20MnCr5) and aluminium (EN AW-6082 T6 condition) is used for the Tailored Forming process chain to produce a load-adapted hollow shaft with high strength areas and reduced weight (see Fig. 1). In order to create a material bond, the semi-finished products with a diameter of $D = 40$ mm are joined by means of rotary friction welding using the parameter set and the procedure described in a previous study [17]. During the process, only the aluminium deforms plastically, as it has a lower strength compared to steel. The resulting aluminium welding flash is removed by machining and the hybrid samples are prepared with a steel length of 22 mm and an aluminium length of 50 mm for the subsequent cup backward extrusion.

<figure><img src="image_762.png" /><figcaption>Figure 1. Process chain of Tailored Forming for a hollow shaft</figcaption></figure>

Prior to the forming process, the hybrid semi-finished workpiece is inductively heated on the steel side in order to create a temperature gradient between steel and aluminium, and consequently an equalisation of the yield stresses. Two heating concepts, with and without cooling on the aluminium side, are used to investigate the influence of the temperature gradient on the joining zone properties. With the shorter component length used here, thermal conduction is even more critical compared to longer sample geometries for full forward extrusion in previous studies [16-18], as only a small volume of steel needs to be heated to a forming temperature of 900 °C. The aluminium with high thermal conductivity must be kept as cool as possible to avoid overheating and melting. As shown in Fig. 2 a) the steel area is heated, while the aluminium is cooled by a sprayed mixture of pressurised air and water from four nozzles [18] using the parameters in Table 1. As the heating of the steel core is only possible by heat conduction due to the skin effect, heating is carried out within 34 s (uncooled) respectively 37 s (cooled) in four stages with decreasing power addition at a frequency of 15.9 kHz to generate a homogeneous temperature
Different extrusion processes were developed and investigated with regard to the bond strength and ductility [16]. Cup backward extrusion proved to be the most suitable process for producing a hybrid hollow shaft due to the greatest influence on the joining zone geometry and consequently the mechanical properties [17]. By using an inductive heating concept with an additional immersion cooling, the temperature gradient between steel and aluminium could be significantly increased in a full forward extrusion process, which led to an improved equalisation of the yield stresses. In this case, the aluminium did not fail in the process, which is why cracks along the joining zone could be largely prevented [16]. Furthermore, in the course of this, an alternative spray cooling was theoretically designed for impact extrusion in the study by Ince et al. [18]. With this concept, greater temperature gradients can be achieved by avoiding the Leidenfrost effect.

Based on the findings presented, the combination of cup backward extrusion combined with an enhanced cooling strategy to achieve a beneficial forming temperature gradient shows high capability to utilise the full potential of the Tailored Forming process chain. This study focuses on cup backward extrusion at elevated forming temperature with partial spray cooling to study the interactions between surface enlargement, the emerging IMC and the resulting interface.

## Materials & Methods

As in previous studies [15], the material pairing of steel (soft-annealed 20MnCr5) and aluminium (EN AW-6082 T6 condition) is used for the Tailored Forming process chain to produce a load-adapted hollow shaft with high strength areas and reduced weight (see Fig. 1). In order to create a material bond, the semi-finished products with a diameter of $D = 40$ mm are joined by means of rotary friction welding using the parameter set and the procedure described in a previous study [17]. During the process, only the aluminium deforms plastically, as it has a lower strength compared to steel. The resulting aluminium welding flash is removed by machining and the hybrid samples are prepared with a steel length of 22 mm and an aluminium length of 50 mm for the subsequent cup backward extrusion.

<figure><img src="image_770.png" /><figcaption>Figure 1. Process chain of Tailored Forming for a hollow shaft</figcaption></figure>

Prior to the forming process, the hybrid semi-finished workpiece is inductively heated on the steel side in order to create a temperature gradient between steel and aluminium, and consequently an equalisation of the yield stresses. Two heating concepts, with and without cooling on the aluminium side, are used to investigate the influence of the temperature gradient on the joining zone properties. With the shorter component length used here, thermal conduction is even more critical compared to longer sample geometries for full forward extrusion in previous studies [16-18], as only a small volume of steel needs to be heated to a forming temperature of 900 °C. The aluminium with high thermal conductivity must be kept as cool as possible to avoid overheating and melting. As shown in Fig. 2 a) the steel area is heated, while the aluminium is cooled by a sprayed mixture of pressurised air and water from four nozzles [18] using the parameters in Table 1. As the heating of the steel core is only possible by heat conduction due to the skin effect, heating is carried out within 34 s (uncooled) respectively 37 s (cooled) in four stages with decreasing power addition at a frequency of 15.9 kHz to generate a homogeneous temperature
distribution. Temperatures were measured using thermocouples type K. To determine the influence of the heat input on the aluminum side, hardness tests are carried out with friction welded samples with and without cooling after induction heating and quenching without forming.

<table><caption>Table 1. Parameters used for spray cooling</caption><thead><tr><th colspan="2">Water</th><th colspan="2">Air</th></tr><tr><th>Pressure in bar</th><th>Mass Flow in l/min</th><th>Pressure in bar</th><th>Mass Flow in m³/h</th></tr></thead><tbody><tr><td>0.7</td><td>1.2</td><td>0.8</td><td>6.8</td></tr></tbody></table>

Cup backward extrusion is carried out on a Lasco SPR 500 screw press. All transfer steps are robot-assisted to ensure reproducibility in the time-sensitive process. At the start of the automated forming process, the sample is inserted into the induction coil by a robot gripper for inductive heating with and without cooling. Once heating is complete, the material is removed automatically from the coil and transferred to the die (see Fig. 2 b)). Subsequently, the punch is pressed into the steel part of the specimen with an impact velocity of 280 mm·s⁻¹. By varying punch diameters between D = 22 mm or D = 30 mm, different cross-sectional changes of 28.8% respectively 53.5% are induced.

<figure><img src="image_763.png" /><figcaption>Figure 2. Experimental set-up: a) induction heating with spray cooling, b) Tool system for cup backward extrusion</figcaption></figure>

After forming, the samples are quenched in water. To avoid thermal cracking, first the aluminium is quenched for 1 s, then the steel for 10 s and the aluminium for 10 s with two repetitions each, before the entire specimen is completely immersed to reach room temperature. The samples are then cut lengthwise to measure the inner and outer diameter D₀, the depth of the joining zone h, the radius r and the width a of the spherical segment (see Fig. 3 a)) to calculate the surface enlargement Aitz. The joining zone consists of a spherical shape with a surface AS and a nearly unformed circular ring AR, which was taken into account as well. The modified surface area Aitz is calculated approximately using the following equation:

$$A_{\text{Iz}} = A_{\text{S}} + A_{\text{R}} = \pi \cdot \left[ (2 \cdot r \cdot h) + \left( \frac{D_{0}^{2} - (2 \cdot a)^{2}}{4} \right) \right] \quad (1)$$
distribution. Temperatures were measured using thermocouples type K. To determine the influence of the heat input on the aluminum side, hardness tests are carried out with friction welded samples with and without cooling after induction heating and quenching without forming.

<table><caption>Table 1. Parameters used for spray cooling</caption><thead><tr><th colspan="2">Water</th><th colspan="2">Air</th></tr><tr><th>Pressure in bar</th><th>Mass Flow in l/min</th><th>Pressure in bar</th><th>Mass Flow in m³/h</th></tr></thead><tbody><tr><td>0.7</td><td>1.2</td><td>0.8</td><td>6.8</td></tr></tbody></table>

Cup backward extrusion is carried out on a Lasco SPR 500 screw press. All transfer steps are robot-assisted to ensure reproducibility in the time-sensitive process. At the start of the automated forming process, the sample is inserted into the induction coil by a robot gripper for inductive heating with and without cooling. Once heating is complete, the material is removed automatically from the coil and transferred to the die (see Fig. 2 b)). Subsequently, the punch is pressed into the steel part of the specimen with an impact velocity of 280 mm·s⁻¹. By varying punch diameters between D = 22 mm or D = 30 mm, different cross-sectional changes of 28.8% respectively 53.5% are induced.

<figure><img src="image_771.png" /><figcaption>Figure 2. Experimental set-up: a) induction heating with spray cooling, b) Tool system for cup backward extrusion</figcaption></figure>

After forming, the samples are quenched in water. To avoid thermal cracking, first the aluminium is quenched for 1 s, then the steel for 10 s and the aluminium for 10 s with two repetitions each, before the entire specimen is completely immersed to reach room temperature. The samples are then cut lengthwise to measure the inner and outer diameter D₀, the depth of the joining zone h, the radius r and the width a of the spherical segment (see Fig. 3 a)) to calculate the surface enlargement Aitz. The joining zone consists of a spherical shape with a surface A₃ and a nearly unformed circular ring Aₘ, which was taken into account as well. The modified surface area Aitz is calculated approximately using the following equation:

$$A_{itz} = A_S + A_R = \pi \cdot \left[ (2 \cdot r \cdot h) + \left( \frac{D_0^2 - (2 \cdot a)^2}{4} \right) \right] \quad (1)$$
Light microscopic and SEM images as well as EDX analyses are performed on the cross sections using a field emission scanning electron microscope Supra VP 55 to evaluate the influence of

cooling effect and the different cross-sectional changes on the joining zone geometry as well as the emerging IMC. The existence of a material bond is examined in different positions (see Fig. 3): in the edge area R0, descending area R1 and R2 as well as in the centre M1 and M0. The local test geometry has a rectangular cross-section of 2.5 mm x 2.5 mm with a length of 8.5 mm and is extracted for each variant at the specified locations using wire eroding. Extracting local areas acts as a test method and allows an evaluation of correlation between surface enlargement, the IMC thickness and the existence of a resulting bond.

<figure><img src="image_764.png" /><figcaption>Figure 3. Geometric quantities and test positions for a material bond</figcaption></figure>

## Results & Discussion

When using spray cooling, there is a significant temperature drop of 314 °C in the aluminium centre $T_{c1}$ (see Fig. 4). In this case, the central steel temperature $T_{c2}$ could even be increased from 800 °C to the necessary 900 °C by adjusting the power profile and heating times, as the aluminium temperatures are well below the melting point due to prevention of excessive heat through the

cooling effect. Therefore, the overall temperature difference can be increased by 116% from 343 °C to 742 °C. The effect of cooling can be clearly seen in the hardness measurements (see Fig. 5). The aluminium base material of the uncooled variant has an average hardness $\bar{x}$ of 77 HV and indicates softening due to excessive heat input. In comparison, cooled samples have an average hardness $\bar{x}_+$ of 92 HV. As long as temperatures of 200 °C are not exceeded for durations of 30 minutes with the alloy EN AW-6082, no strength reductions are to be expected [19]. These two conditions are met in this case, which is why the initial heat treatment condition T6 (97.5 HV10) can be largely maintained by the use of spray cooling. The lower temperatures also have a significant effect on the yield stress in the forming process, which at higher values also require higher forming forces to enable the material to flow (see Fig. 6). The forming resistance of the aluminium

is increased by the cooling during the process and requires a significantly higher forming force of $F = 867$ kN for $D = 30$ mm, which is 38.5% lower without cooling and reaches a value of $F = 626$ kN. A comparison of the punch diameters also shows a difference. As larger volumes have to be formed at $D = 30$ mm, the forming forces required are generally higher with an increase by 104.6% compared to $D = 22$ mm.

<figure><img src="image_765.png" /><figcaption>Figure 4. Temperature-time-curves for cooled and uncooled specimens</figcaption></figure>
Light microscopic and SEM images as well as EDX analyses are performed on the cross sections using a field emission scanning electron microscope Supra VP 55 to evaluate the influence of

cooling effect and the different cross-sectional changes on the joining zone geometry as well as the emerging IMC. The existence of a material bond is examined in different positions (see Fig. 3): in the edge area R0, descending area R1 and R2 as well as in the centre M1 and M0. The local test geometry has a rectangular cross-section of 2.5 mm x 2.5 mm with a length of 8.5 mm and is extracted for each variant at the specified locations using wire eroding. Extracting local areas acts as a test method and allows an evaluation of correlation between surface enlargement, the IMC thickness and the existence of a resulting bond.

<figure><img src="image_772.png" /><figcaption>Figure 3. Geometric quantities and test positions for a material bond</figcaption></figure>

## Results & Discussion

When using spray cooling, there is a significant temperature drop of 314 °C in the aluminium centre $T_{c1}$ (see Fig. 4). In this case, the central steel temperature $T_{c2}$ could even be increased from 800 °C to the necessary 900 °C by adjusting the power profile and heating times, as the aluminium temperatures are well below the melting point due to prevention of excessive heat through the

cooling effect. Therefore, the overall temperature difference can be increased by 116% from 343 °C to 742 °C. The effect of cooling can be clearly seen in the hardness measurements (see Fig. 5). The aluminium base material of the uncooled variant has an average hardness $\bar{x}$ of 77 HV and indicates softening due to excessive heat input. In comparison, cooled samples have an average hardness $\bar{x}_+$ of 92 HV. As long as temperatures of 200 °C are not exceeded for durations of 30 minutes with the alloy EN AW-6082, no strength reductions are to be expected [19]. These two conditions are met in this case, which is why the initial heat treatment condition T6 (97.5 HV10) can be largely maintained by the use of spray cooling. The lower temperatures also have a significant effect on the yield stress in the forming process, which at higher values also require higher forming forces to enable the material to flow (see Fig. 6). The forming resistance of the aluminium

is increased by the cooling during the process and requires a significantly higher forming force of $F = 867$ kN for $D = 30$ mm, which is 38.5% lower without cooling and reaches a value of $F = 626$ kN. A comparison of the punch diameters also shows a difference. As larger volumes have to be formed at $D = 30$ mm, the forming forces required are generally higher with an increase by 104.6% compared to $D = 22$ mm.

<figure><img src="image_773.png" /><figcaption>Figure 4. Temperature-time-curves for cooled and uncooled specimens</figcaption></figure>
<figure><img src="image_766.png" /><figcaption>Figure 5. Hardness curves in the aluminium base material</figcaption></figure>

<figure><img src="image_767.png" /><figcaption>Figure 6. Force-time curves of the cup backward extruded specimens</figcaption></figure>

Depending on the punch diameter and cooling strategy used, different joining zone geometries occur (see Fig. 7). For uncooled variants, the increased temperature of over 450 °C (see Fig. 4) in the aluminium leads to reduced yield strength, which is why the steel area presses into the underlying aluminium during the forming process and gives the joining zone a spherical contour. As a result, a surface enlargement of 52.7% (D = 30 mm, uncooled) is achieved compared to the reference surface (1320.3 mm²), which is only friction welded (see Tab. 2). When cooling is used, the joining zone is less pronounced and is accompanied by a surface enlargement of only 5.9% (D = 30 mm, cooled). The same phenomenon is observed for D = 22 mm, but to a higher extent although a significantly smaller punch cross section (ratio: 53.7%) was used. Reason for this is a higher penetration depth, which, in contrast to D = 30 mm, was guaranteed by sufficient available press energy. The increased cross-section paired with the higher yield stress due to the cooling requires higher forming forces, which were not set in this case. Nevertheless, the cross-sections can be compared with each other on the basis of their percentage surface enlargement.

In the absence of cooling, damage, both in the base material itself and along the IMC is observed, especially in the edge areas a) and g). Depending on the position, IMCs of different thicknesses occur, which reach their maximum in the edge area R0 for all variants. With increasing surface enlargement towards the centre, the IMC becomes thinner. On the one hand, this may be due to the inhomogeneous temperature distribution in the radial direction, which is favoured by
<figure><img src="image_774.png" /><figcaption>Figure 5. Hardness curves in the aluminium base material</figcaption></figure>

<figure><img src="image_775.png" /><figcaption>Figure 6. Force-time curves of the cup backward extruded specimens</figcaption></figure>

Depending on the punch diameter and cooling strategy used, different joining zone geometries occur (see Fig. 7). For uncooled variants, the increased temperature of over 450 °C (see Fig. 4) in the aluminium leads to reduced yield strength, which is why the steel area presses into the underlying aluminium during the forming process and gives the joining zone a spherical contour. As a result, a surface enlargement of 52.7% (D = 30 mm, uncooled) is achieved compared to the reference surface (1320.3 mm²), which is only friction welded (see Tab. 2). When cooling is used, the joining zone is less pronounced and is accompanied by a surface enlargement of only 5.9% (D = 30 mm, cooled). The same phenomenon is observed for D = 22 mm, but to a higher extent although a significantly smaller punch cross section (ratio: 53.7%) was used. Reason for this is a higher penetration depth, which, in contrast to D = 30 mm, was guaranteed by sufficient available press energy. The increased cross-section paired with the higher yield stress due to the cooling requires higher forming forces, which were not set in this case. Nevertheless, the cross-sections can be compared with each other on the basis of their percentage surface enlargement.

In the absence of cooling, damage, both in the base material itself and along the IMC is observed, especially in the edge areas a) and g). Depending on the position, IMCs of different thicknesses occur, which reach their maximum in the edge area R0 for all variants. With increasing surface enlargement towards the centre, the IMC becomes thinner. On the one hand, this may be due to the inhomogeneous temperature distribution in the radial direction, which is favoured by
the skin effect during inductive heating. Higher temperatures occur at the edge, which correlate with thicker IMC. However, the decrease of IMC thickness is intensified by the local increase in surface area towards the core. From studies in which steel and aluminium were joined by compound forging, a decreasing oxide layer content with increasing local plastic strain was also observed [5], which is comparable to the IMC mechanisms occurring here. In the middle, a bond without IMC can be seen in the cooled variants. Two possible scenarios are imaginable for the formation of a bond in the centre:

*   **Scenario 1:** The friction welding process generally results in inferior joint qualities in the centre due to the low relative speeds [13]. In friction welded variants, cracks are present in the centre and are enlarged due to the different thermal expansions of the materials while inductively heated. Subsequent forming counteracts this and creates a material bond in positions c) and i) with IMC and f) and l) without IMC. Due to the significantly greater heat input in the uncooled variants, growth of the IMC is initiated during or after forming. The previous defect is thus closed by the reshaping and an IMC is formed by diffusion.

*   **Scenario 2:** A bond with IMC results from the friction welding process. If an IMC exists, it is either enlarged by inductive heating or breaks apart due to the thermal expansion. If the IMC is intact, it is broken up by the surface enlargement due to the low ductility and flushed outwards by the backward material flow during forming. In outer areas such as e), this leads to an accumulation of the IMC. As a result, juvenile welding surfaces are formed in central areas, which are joined together within forming. This is comparable to the breaking up of oxide layers during cold pressure welding [3]. This phenomenon of IMC rupturing was also observed in a previous study in which punch diameters of *D* = 16 mm were used without cooling [16]. In contrast to larger diameters of 22 mm and 30 mm, however, no IMC was observed for uncooled variants in the centre.

<figure><img src="image_768.png" /><figcaption>Figure 7. Cross sections and SEM images of the different variants with and without cooling</figcaption></figure>

Intermetallic brittle phases such as FeAl₃ or Fe₂Al₅ in particular occur when exposed to heat for a short time. Other types of FeₓAlᵧ such as Fe₃Al or FeAl can be excluded due to the short heating time caused by inductive heating [20]. The observed IMC for uncooled variants exhibits lamellar structures in the positions b) and h). This indicates the formation of the Fe₂Al₅ phase, which emerges at temperatures above 500 °C. By cooling of 500 to 350 °C, the typical rod-shaped phases of the Fe₂Al₅ turn into FeAl₃ [20, 21]. In the edge area, where the IMC is particularly thick, two different phases can be seen both optically by different colouring due to electron backscattering in the SEM image and in the EDX analyses due to different aluminium contents. The steel-side phase has an aluminium content of 68 at.-%, which would indicate the Fe₂Al₅ phase. The aluminium-side phase has a content of 74 at.-% and is indicative of the FeAl₃ phase [20]. In uncooled variants, the formation of only continuous FeAl₃ is suspected due to the lower temperature range of below 350 °C, which is why only one phase can be detected in this case. In the studies by Herbst et
the skin effect during inductive heating. Higher temperatures occur at the edge, which correlate with thicker IMC. However, the decrease of IMC thickness is intensified by the local increase in surface area towards the core. From studies in which steel and aluminium were joined by compound forging, a decreasing oxide layer content with increasing local plastic strain was also observed [5], which is comparable to the IMC mechanisms occurring here. In the middle, a bond without IMC can be seen in the cooled variants. Two possible scenarios are imaginable for the formation of a bond in the centre:

*   **Scenario 1:** The friction welding process generally results in inferior joint qualities in the centre due to the low relative speeds [13]. In friction welded variants, cracks are present in the centre and are enlarged due to the different thermal expansions of the materials while inductively heated. Subsequent forming counteracts this and creates a material bond in positions c) and i) with IMC and f) and l) without IMC. Due to the significantly greater heat input in the uncooled variants, growth of the IMC is initiated during or after forming. The previous defect is thus closed by the reshaping and an IMC is formed by diffusion.

*   **Scenario 2:** A bond with IMC results from the friction welding process. If an IMC exists, it is either enlarged by inductive heating or breaks apart due to the thermal expansion. If the IMC is intact, it is broken up by the surface enlargement due to the low ductility and flushed outwards by the backward material flow during forming. In outer areas such as e), this leads to an accumulation of the IMC. As a result, juvenile welding surfaces are formed in central areas, which are joined together within forming. This is comparable to the breaking up of oxide layers during cold pressure welding [3]. This phenomenon of IMC rupturing was also observed in a previous study in which punch diameters of *D* = 16 mm were used without cooling [16]. In contrast to larger diameters of 22 mm and 30 mm, however, no IMC was observed for uncooled variants in the centre.

<figure><img src="image_776.png" /><figcaption>Figure 7. Cross sections and SEM images of the different variants with and without cooling</figcaption></figure>

Intermetallic brittle phases such as FeAl₃ or Fe₂Al₅ in particular occur when exposed to heat for a short time. Other types of FeₓAlᵧ such as Fe₃Al or FeAl can be excluded due to the short heating time caused by inductive heating [20]. The observed IMC for uncooled variants exhibits lamellar structures in the positions b) and h). This indicates the formation of the Fe₂Al₅ phase, which emerges at temperatures above 500 °C. By cooling of 500 to 350 °C, the typical rod-shaped phases of the Fe₂Al₅ turn into FeAl₃ [20, 21]. In the edge area, where the IMC is particularly thick, two different phases can be seen both optically by different colouring due to electron backscattering in the SEM image and in the EDX analyses due to different aluminium contents. The steel-side phase has an aluminium content of 68 at.-%, which would indicate the Fe₂Al₅ phase. The aluminium-side phase has a content of 74 at.-% and is indicative of the FeAl₃ phase [20]. In uncooled variants, the formation of only continuous FeAl₃ is suspected due to the lower temperature range of below 350 °C, which is why only one phase can be detected in this case. In the studies by Herbst et
al. [14], significant growth of the IMC was only observed at temperatures from 540 °C upwards. As it is only possible to measure the temperatures shortly before forming, therefore a reasonable assumption is, that the aluminium has exceeded this critical temperature due to heat conduction during transfer and especially during forming. The increased contact area of the punch with D = 30 mm in the interface leads to faster diffusion, which is why the IMC thickness b is significantly greater with b = 6.4 µm than at D = 22 mm with b = 3.1 µm. In addition to the use of cooling, the related change in cross-section is therefore another important influencing factor with regard to the IMC formed and consequently the bond strength. In conclusion, the larger contact area and the greater contact pressure during forming causes a stronger growth of the IMC, which can only be prevented by cooling, if larger punch diameters are used.

<table><caption>Table 2. Joining zone surface enlargement ΔA<sub>Jz</sub> and IMC thickness b for different variants</caption><thead><tr><th>D<br>[mm]</th><th>Cooling<br>Strategy</th><th>D<sub>o</sub><br>[mm]</th><th>h<br>[mm]</th><th>a<br>[mm]</th><th>r<br>[mm]</th><th>A<sub>Jz</sub><br>[mm<sup>2</sup>]</th><th>ΔA<sub>Jz</sub><br>[%]</th><th>b<br>[µm]</th></tr></thead><tbody><tr><td>0.0*</td><td>Uncooled</td><td>40.0</td><td>0.0</td><td>0.0</td><td>0.0</td><td>1320.3</td><td>0.0</td><td>n.e.</td></tr><tr><td>0.0*</td><td>Cooled</td><td>40.0</td><td>0.0</td><td>0.0</td><td>0.0</td><td>1320.3</td><td>0.0</td><td>0.7</td></tr><tr><td rowspan="2">22.0</td><td rowspan="2">Uncooled</td><td rowspan="2">41.0</td><td rowspan="2">4.0</td><td rowspan="2">5.1</td><td rowspan="2">14.9</td><td rowspan="2">1709.1</td><td rowspan="2">29.5</td><td>a)</td><td>b)</td><td>c)</td></tr><tr><td>3.1</td><td>2.9</td><td>1.0</td></tr><tr><td rowspan="2">22.0</td><td rowspan="2">Cooled</td><td rowspan="2">41.0</td><td rowspan="2">2.3</td><td rowspan="2">5.5</td><td rowspan="2">23.6</td><td rowspan="2">1470.3</td><td rowspan="2">11.4</td><td>d)</td><td>e)</td><td>f)</td></tr><tr><td>1.0</td><td>0.1</td><td>n.e.</td></tr><tr><td rowspan="2">30.0</td><td rowspan="2">Uncooled</td><td rowspan="2">41.0</td><td rowspan="2">4.8</td><td rowspan="2">6.3</td><td rowspan="2">20.0</td><td rowspan="2">2016.4</td><td rowspan="2">52.7</td><td>g)</td><td>h)</td><td>i)</td></tr><tr><td>6.4</td><td>3.5</td><td>1.0</td></tr><tr><td rowspan="2">30.0</td><td rowspan="2">Cooled</td><td rowspan="2">41.0</td><td rowspan="2">1.6</td><td rowspan="2">6.2</td><td rowspan="2">39.6</td><td rowspan="2">1398.1</td><td rowspan="2">5.9</td><td>j)</td><td>k)</td><td>l)</td></tr><tr><td>1.0</td><td>0.1</td><td>n.e.</td></tr></tbody></table>

*friction welded & not formed

During preparation of rectangular geometries for local bond testing mainly samples without cooling fell apart (see Tab. 3). Friction welded variants (1), (2) without an additional forming step were inductively heated and quenched afterwards to investigate the isolated effect of the heat input on the bond. Due to excessive growth of the IMC, which in turn leads to cracks within the joining zone, the uncooled and unformed sample (1) already fell apart after heating and quenching. Therefore, no bond exists in any position. If cooling is used, the growth of the IMC can be largely suppressed, which is why a bond in position R0, R1, R2 and M1 is present for (2). As previously mentioned, the cooled and unformed sample (2) has a central defect, which is why there is no bond in the position M0. However, this can be compensated for by forming with additional cooling (4), (6).

<table><caption>Table 3. Existence of a material bond depending on position R0, R1, R2, M1, M0</caption><thead><tr><th>Variant</th><th>Position</th><th>R0</th><th>R1</th><th>R2</th><th>M1</th><th>M0</th></tr></thead><tbody><tr><td>(1) Uncooled + Not Formed</td><td></td><td>No Bond</td><td>No Bond</td><td>No Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(2) Cooled + Not Formed</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td></tr><tr><td>(3) 22 Uncooled</td><td></td><td>No Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(4) 22 Cooled</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td></tr><tr><td>(5) 30 Uncooled</td><td></td><td>No Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(6) 30 Cooled</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td></tr></tbody></table>

In the absence of cooling, only specimens in positions R1 and R2 could be extracted of cup backward extruded samples (3), (5). Although, there is a visible bond with an IMC thickness of 1 µm (see Fig. 7) in the centre M0, a premature failure due to small test dimensions occurs. The
al. [14], significant growth of the IMC was only observed at temperatures from 540 °C upwards. As it is only possible to measure the temperatures shortly before forming, therefore a reasonable assumption is, that the aluminium has exceeded this critical temperature due to heat conduction during transfer and especially during forming. The increased contact area of the punch with D = 30 mm in the interface leads to faster diffusion, which is why the IMC thickness b is significantly greater with b = 6.4 µm than at D = 22 mm with b = 3.1 µm. In addition to the use of cooling, the related change in cross-section is therefore another important influencing factor with regard to the IMC formed and consequently the bond strength. In conclusion, the larger contact area and the greater contact pressure during forming causes a stronger growth of the IMC, which can only be prevented by cooling, if larger punch diameters are used.

<table><caption>Table 2. Joining zone surface enlargement ΔA<sub>Jz</sub> and IMC thickness b for different variants</caption><thead><tr><th>D<br>[mm]</th><th>Cooling<br>Strategy</th><th>D<sub>o</sub><br>[mm]</th><th>h<br>[mm]</th><th>a<br>[mm]</th><th>r<br>[mm]</th><th>A<sub>Jz</sub><br>[mm<sup>2</sup>]</th><th>ΔA<sub>Jz</sub><br>[%]</th><th>b<br>[µm]</th></tr></thead><tbody><tr><td>0.0*</td><td>Uncooled</td><td>40.0</td><td>0.0</td><td>0.0</td><td>0.0</td><td>1320.3</td><td>0.0</td><td>n.e.</td></tr><tr><td>0.0*</td><td>Cooled</td><td>40.0</td><td>0.0</td><td>0.0</td><td>0.0</td><td>1320.3</td><td>0.0</td><td>0.7</td></tr><tr><td rowspan="2">22.0</td><td rowspan="2">Uncooled</td><td rowspan="2">41.0</td><td rowspan="2">4.0</td><td rowspan="2">5.1</td><td rowspan="2">14.9</td><td rowspan="2">1709.1</td><td rowspan="2">29.5</td><td>a)</td><td>b)</td><td>c)</td></tr><tr><td>3.1</td><td>2.9</td><td>1.0</td></tr><tr><td rowspan="2">22.0</td><td rowspan="2">Cooled</td><td rowspan="2">41.0</td><td rowspan="2">2.3</td><td rowspan="2">5.5</td><td rowspan="2">23.6</td><td rowspan="2">1470.3</td><td rowspan="2">11.4</td><td>d)</td><td>e)</td><td>f)</td></tr><tr><td>1.0</td><td>0.1</td><td>n.e.</td></tr><tr><td rowspan="2">30.0</td><td rowspan="2">Uncooled</td><td rowspan="2">41.0</td><td rowspan="2">4.8</td><td rowspan="2">6.3</td><td rowspan="2">20.0</td><td rowspan="2">2016.4</td><td rowspan="2">52.7</td><td>g)</td><td>h)</td><td>i)</td></tr><tr><td>6.4</td><td>3.5</td><td>1.0</td></tr><tr><td rowspan="2">30.0</td><td rowspan="2">Cooled</td><td rowspan="2">41.0</td><td rowspan="2">1.6</td><td rowspan="2">6.2</td><td rowspan="2">39.6</td><td rowspan="2">1398.1</td><td rowspan="2">5.9</td><td>j)</td><td>k)</td><td>l)</td></tr><tr><td>1.0</td><td>0.1</td><td>n.e.</td></tr></tbody></table>

*friction welded & not formed

During preparation of rectangular geometries for local bond testing mainly samples without cooling fell apart (see Tab. 3). Friction welded variants (1), (2) without an additional forming step were inductively heated and quenched afterwards to investigate the isolated effect of the heat input on the bond. Due to excessive growth of the IMC, which in turn leads to cracks within the joining zone, the uncooled and unformed sample (1) already fell apart after heating and quenching. Therefore, no bond exists in any position. If cooling is used, the growth of the IMC can be largely suppressed, which is why a bond in position R0, R1, R2 and M1 is present for (2). As previously mentioned, the cooled and unformed sample (2) has a central defect, which is why there is no bond in the position M0. However, this can be compensated for by forming with additional cooling (4), (6).

<table><caption>Table 3. Existence of a material bond depending on position R0, R1, R2, M1, M0</caption><thead><tr><th>Variant</th><th>Position</th><th>R0</th><th>R1</th><th>R2</th><th>M1</th><th>M0</th></tr></thead><tbody><tr><td>(1) Uncooled + Not Formed</td><td></td><td>No Bond</td><td>No Bond</td><td>No Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(2) Cooled + Not Formed</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td></tr><tr><td>(3) 22 Uncooled</td><td></td><td>No Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(4) 22 Cooled</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td></tr><tr><td>(5) 30 Uncooled</td><td></td><td>No Bond</td><td>Bond</td><td>Bond</td><td>No Bond</td><td>No Bond</td></tr><tr><td>(6) 30 Cooled</td><td></td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td><td>Bond</td></tr></tbody></table>

In the absence of cooling, only specimens in positions R1 and R2 could be extracted of cup backward extruded samples (3), (5). Although, there is a visible bond with an IMC thickness of 1 µm (see Fig. 7) in the centre M0, a premature failure due to small test dimensions occurs. The
reason for this is the brittle nature of the very thick IMC, which is occasionally accompanied by cracks in the joining zone regarding the SEM images (see Fig. 7). When cooling is used, at the edge area R0 or in central regions such as M1 or M0 bonds are sufficient enough to withstand the extraction without damage. The positive effects of cooling on the bond are therefore clearly visible and seem to be necessary for defect-free forming.

## Summary & Outlook

In this study, cup backward extrusion of pre-joined hybrid components made of steel and aluminium was investigated with regard to varying punch diameters and heating concepts. Both factors are important control variables in the forming process and have a significant influence on the bond quality due to the sensitive and narrow process windows. A cooling of the aluminium part during the inductive heating process allows the setting of larger temperature gradients and consequently an adjustment of the yield stresses in the forming process. Defects in the aluminium base material that would otherwise occur both outside and inside the joining zone can be prevented. Furthermore, the growth of the IMC is suppressed, which has a positive effect on the local material bond along the interface.

In order to further investigate the influence of the cooling concept and enlargement of the joining zone on the component properties, it is advisable to carry out local tensile tests based on the extracted test cross-sections used in this study. Furthermore, service life tests are necessary to check whether a heat treatment is necessary for sufficient performance. Supplemented by future numerical simulations, a transferability of the correlation between surface enlargement, resulting IMC and increase in bond strength to similar processes within Tailored Forming is ensured.

## Acknowledgments

The results presented in this paper were obtained within sub-project B03 of the Collaborative Research Centre 1153 "Process chain to produce hybrid high performance components with Tailored Forming" funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - SFB 1153 TP B03 - 252662854. The authors thank the DFG for the financial support of this project and the subproject A02 for the SEM/EDX investigations.

## References

[1] ThyssenKrupp Steel AG, "ThyssenKrupp Tailored Blanks bring steel to the right place, compact 26(2) (2007) 32-33.

[2] P. Groche, S. Wohletz, M. Brenneis et al., Joining by forming—A review on joint mechanisms, applications and future trends, Journal of Materials Processing Technology 214(10) (2013) 1972–1994, https://doi.org/10.1016/j.jmatprotec.2013.12.022

[3] N. Bay, Cold Pressure Welding – The Mechanisms Governing Bonding, Journal of Engineering for Industry, (1979) 121-127. https://doi.org/10.1115/1.3439484

[4] P. Groche, S. Wohletz, A. Erbe et al., Effect of the primary heat treatment on the bond formation in cold welding of aluminum and steel, Journal of Materials Processing Technology 214(10) (2014) 2040-2048. https://doi.org/10.1016/j.jmatprotec.2013.12.021

[5] N. J. Peter, C. Gerlitzky, A. Altin et al., Atomic level bonding mechanism in steel/aluminum joints produced by cold pressure welding, Materialia 7 (2019). https://doi.org/10.1016/j.mtla.2019.100396

[6] S. Wohletz and P. Groche, Temperature Influence on Bond Formation in Multi-material Joining by Forging, Procedia Engineering 81 (2014) 2000–2005. https://doi.org/10.1016/j.proeng.2014.10.271

[7] W. Wang, N. J. Politis, Y. Wang et al., Solid-state hot forge bonding of aluminium-steel bimetallic gears: Deformation mechanisms, microstructure and mechanical properties,
reason for this is the brittle nature of the very thick IMC, which is occasionally accompanied by cracks in the joining zone regarding the SEM images (see Fig. 7). When cooling is used, at the edge area R0 or in central regions such as M1 or M0 bonds are sufficient enough to withstand the extraction without damage. The positive effects of cooling on the bond are therefore clearly visible and seem to be necessary for defect-free forming.

## Summary & Outlook

In this study, cup backward extrusion of pre-joined hybrid components made of steel and aluminium was investigated with regard to varying punch diameters and heating concepts. Both factors are important control variables in the forming process and have a significant influence on the bond quality due to the sensitive and narrow process windows. A cooling of the aluminium part during the inductive heating process allows the setting of larger temperature gradients and consequently an adjustment of the yield stresses in the forming process. Defects in the aluminium base material that would otherwise occur both outside and inside the joining zone can be prevented. Furthermore, the growth of the IMC is suppressed, which has a positive effect on the local material bond along the interface.

In order to further investigate the influence of the cooling concept and enlargement of the joining zone on the component properties, it is advisable to carry out local tensile tests based on the extracted test cross-sections used in this study. Furthermore, service life tests are necessary to check whether a heat treatment is necessary for sufficient performance. Supplemented by future numerical simulations, a transferability of the correlation between surface enlargement, resulting IMC and increase in bond strength to similar processes within Tailored Forming is ensured.

## Acknowledgments

The results presented in this paper were obtained within sub-project B03 of the Collaborative Research Centre 1153 "Process chain to produce hybrid high performance components with Tailored Forming" funded by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) - SFB 1153 TP B03 - 252662854. The authors thank the DFG for the financial support of this project and the subproject A02 for the SEM/EDX investigations.

## References

[1] ThyssenKrupp Steel AG, "ThyssenKrupp Tailored Blanks bring steel to the right place, compact 26(2) (2007) 32-33.

[2] P. Groche, S. Wohletz, M. Brenneis et al., Joining by forming—A review on joint mechanisms, applications and future trends, Journal of Materials Processing Technology 214(10) (2013) 1972–1994, https://doi.org/10.1016/j.jmatprotec.2013.12.022

[3] N. Bay, Cold Pressure Welding – The Mechanisms Governing Bonding, Journal of Engineering for Industry, (1979) 121-127. https://doi.org/10.1115/1.3439484

[4] P. Groche, S. Wohletz, A. Erbe et al., Effect of the primary heat treatment on the bond formation in cold welding of aluminum and steel, Journal of Materials Processing Technology 214(10) (2014) 2040-2048. https://doi.org/10.1016/j.jmatprotec.2013.12.021

[5] N. J. Peter, C. Gerlitzky, A. Altin et al., Atomic level bonding mechanism in steel/aluminum joints produced by cold pressure welding, Materialia 7 (2019). https://doi.org/10.1016/j.mtla.2019.100396

[6] S. Wohletz and P. Groche, Temperature Influence on Bond Formation in Multi-material Joining by Forging, Procedia Engineering 81 (2014) 2000–2005. https://doi.org/10.1016/j.proeng.2014.10.271

[7] W. Wang, N. J. Politis, Y. Wang et al., Solid-state hot forge bonding of aluminium-steel bimetallic gears: Deformation mechanisms, microstructure and mechanical properties,
International Journal of Machine Tools & Manufacture 180 (2022).
https://doi.org/10.1016/j.ijmachtools.2022.103930

[8] B.-A. Behrens, A. E. Tekkaya and K.-G. Kosch, Manufacturing of steel-reinforced aluminum parts by co-extrusion and subsequent forging, Key Eng Mater 585 (2014) 149-156.
https://doi.org/10.4028/www.scientific.net/KEM.585.149

[9] A. Foydl, I. Pfeiffer, M. Kammler et al., Manufacturing of Steel-Reinforced Aluminium Products by Combining Hot Extrusion and Closed-Die Forging, Key Engineering Materials, vol. 504-506 (2012) 481-486.
https://doi.org/10.4028/www.scientific.net/KEM.504-506.481

[10] J. Domblesky, F. Kraft, B. Druecke et al., Welded preforms for forging, Journal of Materials Processing Technology 171(1) (2006) 141-149.
https://doi.org/10.1016/j.jmatprotec.2005.06.066

[11] B.-A. Behrens, D. Duran, T. Matthias et al.: Enhancement of the interface friction welded steel-aluminium joints, Prod. Eng. Res. Devel. (2021) 169-176.
https://doi.org/10.1007/s11740-020-00994-5

[12] B.-A. Behrens, J. Uhe; I. Ross et al., Tailored Forming of hybrid bulk metal components, International Journal of Material Forming 15 (2022) #42.
https://doi.org/10.1007/s12289-022-01681-9

[13] H. Ma, G. Qin, Z. Dang and P. Geng, Interfacial microstructure and property of 6061 aluminium alloy/stainless steel hybrid inertia friction welded joint with different steel surface roughness, Materials Characterization. 2021, vol. 179, 111347.
https://doi.org/10.1016/j.matchar.2021.111347

[14] 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.
https://doi.org/10.3139/105.110368

[15] N. Yamamoto, M. Takahashi, M. Aritoshi et al., Formation of intermetallic compounds in friction bonding of Al alloys to steel. Materials Science Forum 539 (2007) 3865-3871.
https://doi.org/10.4028/www.scientific.net/MSF.539-543.3865

[16] A. Piwek, J. Peddinghaus, J. Uhe et al., Influencing the mechanical properties of pre-joined hybrid semi-finished products by impact extrusion, 32nd International Conference on Metallurgy and Materials METAL (2023).
https://doi.org/10.37904/metal.2023.4633

[17] A. Piwek, J. Peddinghaus, J. Uhe et al., Investigation of the joining zone formation of impact extruded components by varied forming sequence and partial cooling, Materials Research Proceedings 28 (2023) 591-600.
https://doi.org/10.21741/9781644902479-64

[18] C.-V. Ince, F. Katz and A. Raatz, Comparative investigation of partial cooling methods for induction heating of hybrid bulk components for hot forming, Materials Research Proceedings 28 (2023) 591-600 (2023) 1083-1090.
https://doi.org/10.21741/9781644902479-64

[19] Gesamtverband der Aluminiumindustrie E.V.: Wärmebehandlung von Aluminiumlegierungen - Merkblatt W7 (2007).

[20] N. Jank, H. Staufer and J. Bruckner, Schweißverbindungen von Stahl mit Aluminium – eine Perspektive für die Zukunft, BHM 153. Jg. Vol. 5 (2008).
https://doi.org/10.1007/s00501-008-0373-6

[21] K. Bouché, F. Barbier and A. Coulet, Intermetallic compound layer growth between solid iron and molten aluminium, Materials Science and Engineering A 249 pp. 167-175 (1998).
https://doi.org/10.1016/S0921-5093(98)00573-5
International Journal of Machine Tools & Manufacture 180 (2022).
https://doi.org/10.1016/j.ijmachtools.2022.103930

[8] B.-A. Behrens, A. E. Tekkaya and K.-G. Kosch, Manufacturing of steel-reinforced aluminum parts by co-extrusion and subsequent forging, Key Eng Mater 585 (2014) 149-156.
https://doi.org/10.4028/www.scientific.net/KEM.585.149

[9] A. Foydl, I. Pfeiffer, M. Kammler et al., Manufacturing of Steel-Reinforced Aluminium Products by Combining Hot Extrusion and Closed-Die Forging, Key Engineering Materials, vol. 504-506 (2012) 481-486.
https://doi.org/10.4028/www.scientific.net/KEM.504-506.481

[10] J. Domblesky, F. Kraft, B. Druecke et al., Welded preforms for forging, Journal of Materials Processing Technology 171(1) (2006) 141-149.
https://doi.org/10.1016/j.jmatprotec.2005.06.066

[11] B.-A. Behrens, D. Duran, T. Matthias et al.: Enhancement of the interface friction welded steel-aluminium joints, Prod. Eng. Res. Devel. (2021) 169-176.
https://doi.org/10.1007/s11740-020-00994-5

[12] B.-A. Behrens, J. Uhe; I. Ross et al., Tailored Forming of hybrid bulk metal components, International Journal of Material Forming 15 (2022) #42.
https://doi.org/10.1007/s12289-022-01681-9

[13] H. Ma, G. Qin, Z. Dang and P. Geng, Interfacial microstructure and property of 6061 aluminium alloy/stainless steel hybrid inertia friction welded joint with different steel surface roughness, Materials Characterization. 2021, vol. 179, 111347.
https://doi.org/10.1016/j.matchar.2021.111347

[14] 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.
https://doi.org/10.3139/105.110368

[15] N. Yamamoto, M. Takahashi, M. Aritoshi et al., Formation of intermetallic compounds in friction bonding of Al alloys to steel. Materials Science Forum 539 (2007) 3865-3871.
https://doi.org/10.4028/www.scientific.net/MSF.539-543.3865

[16] A. Piwek, J. Peddinghaus, J. Uhe et al., Influencing the mechanical properties of pre-joined hybrid semi-finished products by impact extrusion, 32nd International Conference on Metallurgy and Materials METAL (2023).
https://doi.org/10.37904/metal.2023.4633

[17] A. Piwek, J. Peddinghaus, J. Uhe et al., Investigation of the joining zone formation of impact extruded components by varied forming sequence and partial cooling, Materials Research Proceedings 28 (2023) 591-600.
https://doi.org/10.21741/9781644902479-64

[18] C.-V. Ince, F. Katz and A. Raatz, Comparative investigation of partial cooling methods for induction heating of hybrid bulk components for hot forming, Materials Research Proceedings 28 (2023) 591-600 (2023) 1083-1090.
https://doi.org/10.21741/9781644902479-64

[19] Gesamtverband der Aluminiumindustrie E.V.: Wärmebehandlung von Aluminiumlegierungen - Merkblatt W7 (2007).

[20] N. Jank, H. Staufer and J. Bruckner, Schweißverbindungen von Stahl mit Aluminium – eine Perspektive für die Zukunft, BHM 153. Jg. Vol. 5 (2008).
https://doi.org/10.1007/s00501-008-0373-6

[21] K. Bouché, F. Barbier and A. Coulet, Intermetallic compound layer growth between solid iron and molten aluminium, Materials Science and Engineering A 249 pp. 167-175 (1998).
https://doi.org/10.1016/S0921-5093(98)00573-5
# Using computer vision to analyse fracture strains of oxide scale layers on a macro level

WESTER Hendrik¹,ᵃ*, HUNZE-TRETOW Jan Niklas¹,ᵇ, BRUNOTTE Kai¹,ᶜ and BEHRENS Bernd-Arno¹,ᵈ

¹Institute for Forming Technology and Forming Machines, An der Universität 2, 30823 Garbsen, Germany

**Keywords:** Computer Vision, Fracture Types, Oxide Scale, Tensile Test

**Abstract.** Hot forging has established itself as an efficient process for the manufacture of highly stressed components. The high semi-finished product temperatures significantly increase the deformation capacity and enable the production of complex geometries. However, high semi-finished product temperatures of up to 1250 °C also lead to increased oxide scale formation. Therefore, oxide scale plays an important role in the context of hot forming processes. Due to the contrasting properties between steel substrates and oxide scale, the appearance of oxide scale affects numerous influencing factors, such as changed friction conditions or thermophysical properties. With increasing interest in numerical process prediction arises the demand to take into account the behaviour of oxide scale in finite-element simulations. In addition to the numerical mapping of the crack behaviour, the challenge in mapping the oxide scale is to determine suitable parameters for describing the failure behaviour. Therefore, this work focuses on a novel procedure to characterise the failure of oxide scale under process relevant conditions of hot forging.

**Introduction**

Hot forging is widely used to manufacture industrial components. The pre-heating of semi-finished products to temperatures between 800 °C and 1250 °C, significantly decrease the required forming forces and increase formability. However, the high temperatures lead to the formation of oxide scale [1]. Depending on the steel matrix properties like carbon content as well as temperature and oxidation time, different iron oxides like wustite, haematite and magnetite occur. The influence of alloying elements on the oxidation behaviour of steel has been described in [2]. The oxide scale has to be removed which lead to material losses. In addition, the oxide scale particularly influences friction as well as material flow and can lead to an increase of die wear [3]. Furthermore, it was stated that an oxide scale layer can have a protective character. However, protection only takes place if the oxide scale layer is free from defects like cracks or spallation zones [4]. Within hot forming applications, oxide scale is temporarily categorised into primary, secondary and tertiary oxide scale. These different types are formed under different process conditions and react differently to external stresses during the process. They also differ in terms of their volumetric formation. The primary oxide scale is formed during reheating of the material, which usually takes place in conventional furnaces. In case of a slow heating of the material within a furnace before the forming process this type of oxide scale can form unhindered. Therefore, primary oxide scale is also the most pronounced in terms of mass. The primary oxide scale is ideally removed by a high-pressure scale washer or by pre-dipping the samples from the surface, which results in a high material loss. As oxide scale grows immediately in atmosphere, a new thin oxide scale layer develops on the descaled steel surface [5]. Thus, during hot forging in particular the secondary oxide scale has to be considered, which is formed during the transfer into the forming press or between forming steps. Due to shorter oxidation times the layer thickness is small compared to

<figure><img src="image_769.png" /></figure>

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

<footnote>ᵃwester@ifum.uni-hannover.de, ᵇhunze@ifum.uni-hannover.de,</footnote>

<footnote>ᶜbrunotte@ifum.uni-hannover.de, ᵈbehrens@ifum.uni-hannover.de</footnote>
# Using computer vision to analyse fracture strains of oxide scale layers on a macro level

WESTER Hendrik¹,ᵃ*, HUNZE-TRETOW Jan Niklas¹,ᵇ, BRUNOTTE Kai¹,ᶜ and BEHRENS Bernd-Arno¹,ᵈ

¹Institute for Forming Technology and Forming Machines, An der Universität 2, 30823 Garbsen, Germany

**Keywords:** Computer Vision, Fracture Types, Oxide Scale, Tensile Test

**Abstract.** Hot forging has established itself as an efficient process for the manufacture of highly stressed components. The high semi-finished product temperatures significantly increase the deformation capacity and enable the production of complex geometries. However, high semi-finished product temperatures of up to 1250 °C also lead to increased oxide scale formation. Therefore, oxide scale plays an important role in the context of hot forming processes. Due to the contrasting properties between steel substrates and oxide scale, the appearance of oxide scale affects numerous influencing factors, such as changed friction conditions or thermophysical properties. With increasing interest in numerical process prediction arises the demand to take into account the behaviour of oxide scale in finite-element simulations. In addition to the numerical mapping of the crack behaviour, the challenge in mapping the oxide scale is to determine suitable parameters for describing the failure behaviour. Therefore, this work focuses on a novel procedure to characterise the failure of oxide scale under process relevant conditions of hot forging.

**Introduction**

Hot forging is widely used to manufacture industrial components. The pre-heating of semi-finished products to temperatures between 800 °C and 1250 °C, significantly decrease the required forming forces and increase formability. However, the high temperatures lead to the formation of oxide scale [1]. Depending on the steel matrix properties like carbon content as well as temperature and oxidation time, different iron oxides like wustite, haematite and magnetite occur. The influence of alloying elements on the oxidation behaviour of steel has been described in [2]. The oxide scale has to be removed which lead to material losses. In addition, the oxide scale particularly influences friction as well as material flow and can lead to an increase of die wear [3]. Furthermore, it was stated that an oxide scale layer can have a protective character. However, protection only takes place if the oxide scale layer is free from defects like cracks or spallation zones [4]. Within hot forming applications, oxide scale is temporarily categorised into primary, secondary and tertiary oxide scale. These different types are formed under different process conditions and react differently to external stresses during the process. They also differ in terms of their volumetric formation. The primary oxide scale is formed during reheating of the material, which usually takes place in conventional furnaces. In case of a slow heating of the material within a furnace before the forming process this type of oxide scale can form unhindered. Therefore, primary oxide scale is also the most pronounced in terms of mass. The primary oxide scale is ideally removed by a high-pressure scale washer or by pre-dipping the samples from the surface, which results in a high material loss. As oxide scale grows immediately in atmosphere, a new thin oxide scale layer develops on the descaled steel surface [5]. Thus, during hot forging in particular the secondary oxide scale has to be considered, which is formed during the transfer into the forming press or between forming steps. Due to shorter oxidation times the layer thickness is small compared to

<figure><img src="image_777.png" /></figure>

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

<footnote>ᵃwester@ifum.uni-hannover.de, ᵇhunze@ifum.uni-hannover.de,</footnote>

<footnote>ᶜbrunotte@ifum.uni-hannover.de, ᵈbehrens@ifum.uni-hannover.de</footnote>
primary oxide scale. In contrast to thick layers of scale that break up and crack, thin scale deforms evenly with the steel matrix during hot rolling [5]. Based on process conditions like oxidation time and temperature layer thicknesses between 10 µm and 140 µm are formed. Higher oxidation temperatures as well as longer oxidation times leads to an increased oxide scale formation [6]. The development of the oxide scale layer thickness as a function of temperature and oxidation time as well as the alloying elements carbon and chromium was investigated in detail by [6]. Based on the experimental results, a calculation approach was developed to predict the oxide scale layer thickness.

It is known that oxide scale exhibits different deformation behaviour depending on its composition and the existing process conditions. Experimental investigations at elevated temperatures within a range between 900 °C and 1200 °C have shown the different material-specific deformation behaviour. The highest bearable strain was observed for wustite whereas its yield stress has been found to be lower as compared to haematite and magnetite. Haematite has been observed to possess the highest yield stress. Moreover, it has also been found to be the hardest oxide scale component at room temperature with a Vickers hardness of 1000 HV10 as compared to magnetite (600 HV10) and wustite (400 HV10). Furthermore, investigations have shown a significant influence of the strain rate on the forming behaviour [7, 8]. The deformation behaviour of oxidised low carbon steel samples was described in [9]. It was shown, that each phase of the oxide scale changes its deformation behaviour from elastic brittle to plastic at a specific temperature. A strong dependence of the deformation behaviour on the test temperature was also observed during plane strain compression tests of oxidised steel samples under hot rolling conditions. With decreasing test temperature the achievable deformation of oxide scale is reduced [10].

As the oxide scale is hard and brittle compared to the steel substrate it often cracks and separates during forming operations. In [11] diagrams are presented that can be used to predict different types of oxide scale failure, which can occur under tensile or compressive conditions. It was shown, that depending on the stress conditions that acted on oxide scale, phenomena like delamination, crack formation or spallation could be predicted. The presented diagrams and mathematical approaches are a proper tool to predict the type of failure and can be used as basis for numerical finite element simulations to model oxide scale failure. For realistic modelling, however, specific material data is required, for which suitable test methods are needed that can reproduce process-relevant conditions. In this context, the fracture properties of oxide scale under hot rolling conditions were studied. The experimental tests showed, that oxide scale failure mode is strongly dependent on the temperature and the composition of the oxide scale [12]. The oxide scale fracture was investigated by tensile tests using round tensile specimens. The tests were performed under variation of forming speed and forming temperature. Prior to forming, the samples were subjected to defined oxidation in an air atmosphere. The obtained results indicate two failure modes leading to oxide spallation which are strongly influenced by the temperature, strain rate and strain. The first mode occurs at low temperatures between 850 °C – 900 °C and is characterised by an initial through cracking. This initial cracking is followed by initiation and propagation of a crack along the oxide-metal interface between adjacent cracks. A second failure mode was observed at higher temperature between 900 °C – 1200 °C. As a slipping of the oxide scale relative to the metal surface and ductile behaviour of it has been observed for this mode, it indicated a weak interface and strong oxide scale.

The test methodology presented in [12] provided fundamental findings on the damage of oxide scale layers under conditions relevant to hot rolling. However, the methodology also poses some challenges. For example, it is not possible to clearly determine the time of crack initiation. But, knowledge of the crack initiation time is of particular interest for determining material parameters to describe the damage behaviour such as critical strains. This paper therefore presents a new test
primary oxide scale. In contrast to thick layers of scale that break up and crack, thin scale deforms evenly with the steel matrix during hot rolling [5]. Based on process conditions like oxidation time and temperature layer thicknesses between 10 µm and 140 µm are formed. Higher oxidation temperatures as well as longer oxidation times leads to an increased oxide scale formation [6]. The development of the oxide scale layer thickness as a function of temperature and oxidation time as well as the alloying elements carbon and chromium was investigated in detail by [6]. Based on the experimental results, a calculation approach was developed to predict the oxide scale layer thickness.

It is known that oxide scale exhibits different deformation behaviour depending on its composition and the existing process conditions. Experimental investigations at elevated temperatures within a range between 900 °C and 1200 °C have shown the different material-specific deformation behaviour. The highest bearable strain was observed for wustite whereas its yield stress has been found to be lower as compared to haematite and magnetite. Haematite has been observed to possess the highest yield stress. Moreover, it has also been found to be the hardest oxide scale component at room temperature with a Vickers hardness of 1000 HV10 as compared to magnetite (600 HV10) and wustite (400 HV10). Furthermore, investigations have shown a significant influence of the strain rate on the forming behaviour [7, 8]. The deformation behaviour of oxidised low carbon steel samples was described in [9]. It was shown, that each phase of the oxide scale changes its deformation behaviour from elastic brittle to plastic at a specific temperature. A strong dependence of the deformation behaviour on the test temperature was also observed during plane strain compression tests of oxidised steel samples under hot rolling conditions. With decreasing test temperature the achievable deformation of oxide scale is reduced [10].

As the oxide scale is hard and brittle compared to the steel substrate it often cracks and separates during forming operations. In [11] diagrams are presented that can be used to predict different types of oxide scale failure, which can occur under tensile or compressive conditions. It was shown, that depending on the stress conditions that acted on oxide scale, phenomena like delamination, crack formation or spallation could be predicted. The presented diagrams and mathematical approaches are a proper tool to predict the type of failure and can be used as basis for numerical finite element simulations to model oxide scale failure. For realistic modelling, however, specific material data is required, for which suitable test methods are needed that can reproduce process-relevant conditions. In this context, the fracture properties of oxide scale under hot rolling conditions were studied. The experimental tests showed, that oxide scale failure mode is strongly dependent on the temperature and the composition of the oxide scale [12]. The oxide scale fracture was investigated by tensile tests using round tensile specimens. The tests were performed under variation of forming speed and forming temperature. Prior to forming, the samples were subjected to defined oxidation in an air atmosphere. The obtained results indicate two failure modes leading to oxide spallation which are strongly influenced by the temperature, strain rate and strain. The first mode occurs at low temperatures between 850 °C – 900 °C and is characterised by an initial through cracking. This initial cracking is followed by initiation and propagation of a crack along the oxide-metal interface between adjacent cracks. A second failure mode was observed at higher temperature between 900 °C – 1200 °C. As a slipping of the oxide scale relative to the metal surface and ductile behaviour of it has been observed for this mode, it indicated a weak interface and strong oxide scale.

The test methodology presented in [12] provided fundamental findings on the damage of oxide scale layers under conditions relevant to hot rolling. However, the methodology also poses some challenges. For example, it is not possible to clearly determine the time of crack initiation. But, knowledge of the crack initiation time is of particular interest for determining material parameters to describe the damage behaviour such as critical strains. This paper therefore presents a new test
and evaluation methodology to investigate the damage behaviour of oxide scale layers in detail under process-relevant conditions for forging. The data obtained should enable a numerical representation of the damage behaviour of scale in future work.

## Materials and Methods

**Experimental test procedure.** Based on the procedure presented in [12], a tensile test is used to analyse the oxide scale damage behaviour. As described, the focus is on a methodology for determining the time of crack initiation. In addition, it should also be possible to precisely analyse the type of damage. For this purpose, the tensile test is used together with an optical 2D measuring system GOM ARAMIS. By using an optical measuring system, the test can be analysed with both temporal and local resolution. Flat tensile specimens are used in order to be able to analyse as large an area of the specimen as possible. In contrast to the round tensile specimens used in [12], the entire upper side of the specimen can be analysed with flat tensile specimens. The test setup is shown schematically in Fig. 1 (a). A DIL805 A/D+T quenching and forming dilatometer is used in tensile mode for the tensile tests. The tensile specimen is fixed in the left and right specimen holders and then inductively heated using a flat coil. The temperature is controlled via welded-on type S thermocouples. The sample geometry used for tensile tests within forming dilatometer is shown in Fig. 1 (b).

<figure><img src="image_770.png" /><figcaption>Figure 1 Schematic test setup (a) and sample geometry (b)</figcaption></figure>

The measuring range of the specimen is 10 mm long and 3 mm wide. The thickness of the samples was 1.5 mm. C45 steel was chosen as the material as knowledge on oxide scale growth is already available for this material from previous work. Furthermore, it only consists of the alloying element carbon, so that the influence of carbon on oxide scale damage can be considered later by looking at other alloys such as C15 or C60. During the whole test procedure which includes heating, oxidisation and mechanical testing the sample is fixed between the sample holders within the test chamber. Before the mechanical test under tensile load, the specimens are oxidised in a targeted manner. For this purpose, the samples are first inductively heated with a heating rate of 10 K/s to a specified oxidation temperature under vacuum in the test chamber. After reaching the test temperature, an air atmosphere is set in the test chamber and the specimen is specifically oxidised for a defined time. Once the selected oxidation time has been reached, the test chamber is flooded with nitrogen to prevent further oxidation. For all tests an oxidation time of 120 s were used. After oxidation, the sample is mechanical deformed at a defined test temperature up to a degree of deformation of 0.2. First, the test temperature is chosen to be the same as the oxidation temperature. In further tests, the sample is cooled below the test temperature before the mechanical test. Cooling was carried out at a cooling rate of 5 K/s from oxidation temperature to test temperature using nitrogen. Temperatures were controlled using a thermocouple Type S welded onto the sample surface. A schematic test scheme is shown in Fig. 2. Each test is recorded with the optical measuring system so that time-resolved image series are subsequently available for
and evaluation methodology to investigate the damage behaviour of oxide scale layers in detail under process-relevant conditions for forging. The data obtained should enable a numerical representation of the damage behaviour of scale in future work.

## Materials and Methods

**Experimental test procedure.** Based on the procedure presented in [12], a tensile test is used to analyse the oxide scale damage behaviour. As described, the focus is on a methodology for determining the time of crack initiation. In addition, it should also be possible to precisely analyse the type of damage. For this purpose, the tensile test is used together with an optical 2D measuring system GOM ARAMIS. By using an optical measuring system, the test can be analysed with both temporal and local resolution. Flat tensile specimens are used in order to be able to analyse as large an area of the specimen as possible. In contrast to the round tensile specimens used in [12], the entire upper side of the specimen can be analysed with flat tensile specimens. The test setup is shown schematically in Fig. 1 (a). A DIL805 A/D+T quenching and forming dilatometer is used in tensile mode for the tensile tests. The tensile specimen is fixed in the left and right specimen holders and then inductively heated using a flat coil. The temperature is controlled via welded-on type S thermocouples. The sample geometry used for tensile tests within forming dilatometer is shown in Fig. 1 (b).

<figure><img src="image_778.png" /><figcaption>Figure 1 Schematic test setup (a) and sample geometry (b)</figcaption></figure>

The measuring range of the specimen is 10 mm long and 3 mm wide. The thickness of the samples was 1.5 mm. C45 steel was chosen as the material as knowledge on oxide scale growth is already available for this material from previous work. Furthermore, it only consists of the alloying element carbon, so that the influence of carbon on oxide scale damage can be considered later by looking at other alloys such as C15 or C60. During the whole test procedure which includes heating, oxidisation and mechanical testing the sample is fixed between the sample holders within the test chamber. Before the mechanical test under tensile load, the specimens are oxidised in a targeted manner. For this purpose, the samples are first inductively heated with a heating rate of 10 K/s to a specified oxidation temperature under vacuum in the test chamber. After reaching the test temperature, an air atmosphere is set in the test chamber and the specimen is specifically oxidised for a defined time. Once the selected oxidation time has been reached, the test chamber is flooded with nitrogen to prevent further oxidation. For all tests an oxidation time of 120 s were used. After oxidation, the sample is mechanical deformed at a defined test temperature up to a degree of deformation of 0.2. First, the test temperature is chosen to be the same as the oxidation temperature. In further tests, the sample is cooled below the test temperature before the mechanical test. Cooling was carried out at a cooling rate of 5 K/s from oxidation temperature to test temperature using nitrogen. Temperatures were controlled using a thermocouple Type S welded onto the sample surface. A schematic test scheme is shown in Fig. 2. Each test is recorded with the optical measuring system so that time-resolved image series are subsequently available for
analysing the oxide scale damage behaviour. The exposure time setting is of particular importance for the optical analysis of the damage behaviour. This determines the brightness of the respective image. Depending on the test temperature, different settings are required for the exposure time. The maximum possible exposure frequency also goes hand in hand with the exposure time. In principle, exposure frequency should be set as high as possible in order to generate the maximum information content. An exposure time of 1 ms is selected at 1100 °C. This allows the camera's maximum exposure frequency of 160 Hz to be used. At the lowest tested temperature, the exposure time needs to increased up to 13 ms, which leads to an exposure frequency of 65 Hz.

<figure><img src="image_771.png" /><figcaption>Figure 2 Schematic Test setup</figcaption></figure>

Within this study the influence of different oxidation and test temperatures as well as strain rates were investigated. An overview over the varied parameters, which should represent the hot forging-relevant process conditions, is given in Table 1. A lower test temperature than the oxidation temperature is intended to represent a possible cooling of the oxide scale in air or through contact with the cold tools. Due to the large combination of parameters, the test matrix was not initially subjected to a full factorial test. When selecting the parameters, the focus was on the individual consideration of various effects such as the influence of the oxidation temperature, the influence of cooling after oxidation or the influence of the strain rate. The tested parameter combinations are discussed in the results section. Each test was repeated 5 times for statistical validation

<table><caption>Table 1. Test Matrix</caption><tr><td>Oxidation temperature [°C]</td><td>1100, 1050, 1000, 950, 900, 850</td></tr><tr><td>Test temperature [°C]</td><td>1100, 1050, 1000, 950, 900, 750, 600, 450</td></tr><tr><td>Strain rate [s<sup>-1</sup>]</td><td>0.01, 0.05, 0.1</td></tr><tr><td>Oxidation time [s]</td><td>120</td></tr></table>

**Evaluation Method:** Following the test, the Aramis system delivers an image sequence. Each image is saved in a separate file in greyscale pixel format with a size of 2400 pixel in x direction and 1728 pixel in y direction. This means that there is a two-dimensional arrangement of values, which can be integers between zero (black) and 255 (white). As large amounts of data are generated in the course of the test evaluation, a semi-automated Python script was developed using SciPy library for data processing based on computer vision approaches.

First, the crack initiation point is determined automatically. The images are called up one after the other in chronological order and analysed for their characteristics. To speed up this iteration, not the consecutive images are called up, but every 40th image in a first loop. If a crack is found,
analysing the oxide scale damage behaviour. The exposure time setting is of particular importance for the optical analysis of the damage behaviour. This determines the brightness of the respective image. Depending on the test temperature, different settings are required for the exposure time. The maximum possible exposure frequency also goes hand in hand with the exposure time. In principle, exposure frequency should be set as high as possible in order to generate the maximum information content. An exposure time of 1 ms is selected at 1100 °C. This allows the camera's maximum exposure frequency of 160 Hz to be used. At the lowest tested temperature, the exposure time needs to increased up to 13 ms, which leads to an exposure frequency of 65 Hz.

<figure><img src="image_779.png" /><figcaption>Figure 2 Schematic Test setup</figcaption></figure>

Within this study the influence of different oxidation and test temperatures as well as strain rates were investigated. An overview over the varied parameters, which should represent the hot forging-relevant process conditions, is given in Table 1. A lower test temperature than the oxidation temperature is intended to represent a possible cooling of the oxide scale in air or through contact with the cold tools. Due to the large combination of parameters, the test matrix was not initially subjected to a full factorial test. When selecting the parameters, the focus was on the individual consideration of various effects such as the influence of the oxidation temperature, the influence of cooling after oxidation or the influence of the strain rate. The tested parameter combinations are discussed in the results section. Each test was repeated 5 times for statistical validation

<table><caption>Table 1. Test Matrix</caption><tr><td>Oxidation temperature [°C]</td><td>1100, 1050, 1000, 950, 900, 850</td></tr><tr><td>Test temperature [°C]</td><td>1100, 1050, 1000, 950, 900, 750, 600, 450</td></tr><tr><td>Strain rate [s<sup>-1</sup>]</td><td>0.01, 0.05, 0.1</td></tr><tr><td>Oxidation time [s]</td><td>120</td></tr></table>

**Evaluation Method:** Following the test, the Aramis system delivers an image sequence. Each image is saved in a separate file in greyscale pixel format with a size of 2400 pixel in x direction and 1728 pixel in y direction. This means that there is a two-dimensional arrangement of values, which can be integers between zero (black) and 255 (white). As large amounts of data are generated in the course of the test evaluation, a semi-automated Python script was developed using SciPy library for data processing based on computer vision approaches.

First, the crack initiation point is determined automatically. The images are called up one after the other in chronological order and analysed for their characteristics. To speed up this iteration, not the consecutive images are called up, but every 40th image in a first loop. If a crack is found,
a more precise search is carried out in two stages between the chosen image and the previously analysed image. For crack identification, a search area is first defined, the so-called crack identification box (CIB). This is done for the initial image of the image series. In this image, a rectangular area is defined which covers the gauge area of the sample. However, the CIB should be smaller than the gauge area of the sample in order not to detect localised changes at the sample edges. To determine the CIB, a horizontal Sobel filter is applied to the image array to emphasize horizontal edges. This creates a gradient image from the original image. The areas of highest intensity are where the brightness of the original image changes the most and therefore represent the sharpest edges [13]. Depending on the orientation, the edges take on a very small or very large value in the array of the image. The Python script utilises this property and extracts the gauge area using defined limit values. Fig. 3 (a) shows the processed image and the defined CIB. The generated CIB can also be recognised as a superimposed rectangle and the vertical line, which is used to identify the horizontal sample limits. An initial analysis of the tests showed that, for test temperatures between 849 °C and 1051 °C ductile cracks can be observed. Outside this temperature range brittle cracks occur. Furthermore, a test temperature below 849 °C thermal shifts due to cooling and at temperatures above 1051 °C formation of bubbles can be observed. To consider the different failure behaviour and the described phenomena two methods were used for crack identification. For temperatures within the defined range, where ductile cracks can be observed, a mean value kernel with a size of (20,1) and a stride or step size of the kernel of (1,5) is used. Here, the size of the kernel describes the size of the mask to which the filter is applied and the stride the shift of this mask. The lines of the convolved image are then considered and a Savitzky-Golay filter is applied to the individual lines to reduce noise [14, 15]. Subsequently the gradient is calculated. A crack is characterised by the fact that a local maximum follows a local minimum in a small interval, or vice versa. Because cracks are known to propagate in a vertical direction, the mean value filter with its vertical shape can be used to reduce the noise. Due to its shape, a crack will exhibit a characteristic value in the filtered matrix despite the filter (Fig. 3 (b)).

<figure><img src="image_772.png" /><figcaption>Figure 3 Definition of Crack Identification Box based on the first image (a) and processed image for identification of ductile cracks within the oxide scale (b)</figcaption></figure>

If the test temperature is outside an interval of 849 °C and 1051 °C, a different crack detection is performed. This is necessary because phenomena like bubbles and thermal shifts can lead to the detection of vertical edges on the sample, which can be incorrectly interpreted as cracks. However, as this is a static occurrence, the same information content is present on image *i* and on a previous image. Therefore, in addition to the current image (*i*), the penultimate image (*i*-2) is also used. The two images are first filtered for vertical edges using a Sobel filter. This is shown in Fig. 4 (a) for image *i*. The two filtered images are then subtracted from each other to determine a total gradient of the vertical edges. The static effects are filtered out by subtraction like shown in Fig. 4 (b). As
a more precise search is carried out in two stages between the chosen image and the previously analysed image. For crack identification, a search area is first defined, the so-called crack identification box (CIB). This is done for the initial image of the image series. In this image, a rectangular area is defined which covers the gauge area of the sample. However, the CIB should be smaller than the gauge area of the sample in order not to detect localised changes at the sample edges. To determine the CIB, a horizontal Sobel filter is applied to the image array to emphasize horizontal edges. This creates a gradient image from the original image. The areas of highest intensity are where the brightness of the original image changes the most and therefore represent the sharpest edges [13]. Depending on the orientation, the edges take on a very small or very large value in the array of the image. The Python script utilises this property and extracts the gauge area using defined limit values. Fig. 3 (a) shows the processed image and the defined CIB. The generated CIB can also be recognised as a superimposed rectangle and the vertical line, which is used to identify the horizontal sample limits. An initial analysis of the tests showed that, for test temperatures between 849 °C and 1051 °C ductile cracks can be observed. Outside this temperature range brittle cracks occur. Furthermore, a test temperature below 849 °C thermal shifts due to cooling and at temperatures above 1051 °C formation of bubbles can be observed. To consider the different failure behaviour and the described phenomena two methods were used for crack identification. For temperatures within the defined range, where ductile cracks can be observed, a mean value kernel with a size of (20,1) and a stride or step size of the kernel of (1,5) is used. Here, the size of the kernel describes the size of the mask to which the filter is applied and the stride the shift of this mask. The lines of the convolved image are then considered and a Savitzky-Golay filter is applied to the individual lines to reduce noise [14, 15]. Subsequently the gradient is calculated. A crack is characterised by the fact that a local maximum follows a local minimum in a small interval, or vice versa. Because cracks are known to propagate in a vertical direction, the mean value filter with its vertical shape can be used to reduce the noise. Due to its shape, a crack will exhibit a characteristic value in the filtered matrix despite the filter (Fig. 3 (b)).

<figure><img src="image_780.png" /><figcaption>Figure 3 Definition of Crack Identification Box based on the first image (a) and processed image for identification of ductile cracks within the oxide scale (b)</figcaption></figure>

If the test temperature is outside an interval of 849 °C and 1051 °C, a different crack detection is performed. This is necessary because phenomena like bubbles and thermal shifts can lead to the detection of vertical edges on the sample, which can be incorrectly interpreted as cracks. However, as this is a static occurrence, the same information content is present on image *i* and on a previous image. Therefore, in addition to the current image (*i*), the penultimate image (*i*-2) is also used. The two images are first filtered for vertical edges using a Sobel filter. This is shown in Fig. 4 (a) for image *i*. The two filtered images are then subtracted from each other to determine a total gradient of the vertical edges. The static effects are filtered out by subtraction like shown in Fig. 4 (b). As
the cracks in this temperature range are "brittle" cracks that propagate rapidly, they can be identified using this method with a step size of 2. Using this processed image based in image *i* and image *i-2* the same operations are used for crack identification as for ductile cracks. However, the kernel has the size (160,1) and takes place in a stride of (1,5). As brittle cracks are more pronounced, they can be better detected using the method described. A larger kernel can therefore be used. As with "ductile" cracks, the shape of the crack is used for identification. Cracks are searched for along the horizontal. To reduce the noise, the Savitzky-Golay filter is applied again. The final identification of a crack is carried out analogously to the ductile cracks by the sequence of minima and maxima. To avoid errors, the identified cracks are checked and, if necessary, adjusted by the user as part of the evaluation process.

<figure><img src="image_773.png" /><figcaption>Figure 4 Processed images to identify brittle cracks based on image *i* (a) and total gradient image using image *i* and *i-2* (b)</figcaption></figure>

Finally, the sample is measured optically. The first image, i.e. the image in which there is no deformation, and the image that represents the time of crack initiation are used. The measurement is based on the distance between the shoulders of the sample. For this purpose, a Sobel filter is applied to both images to emphasise vertical edges. An average kernel of size (40,1) is subsequently applied to both filtered images along a line to take into account the shape of the shoulder as a vertical line and reduce noise. The edges of the shoulder can be identified via the maximum or minimum of the array starting from the beginning or end. Fig. 5 shows an example of this for the "brittle" crack. The edges of the sample shoulders are clearly recognisable in all images and were correctly identified by the algorithm. Finally, the strain is calculated using the measured sample lengths. The sample length itself is taken as a unitless proportion of the total image width. A critical strain which leads to failure of the oxide scale layer is calculated as the change in length ($l_1-l_0$) of its original length ($l_0$). As incorrect crack detection cannot be completely prevented, a final check is carried out by manual user control of the identified crack time and the determined line for measuring the change in length. For this purpose, the corresponding images are displayed after the evaluation. This was carried out for all evaluated tests, with the algorithm demonstrating a very high identification rate.
the cracks in this temperature range are "brittle" cracks that propagate rapidly, they can be identified using this method with a step size of 2. Using this processed image based in image *i* and image *i-2* the same operations are used for crack identification as for ductile cracks. However, the kernel has the size (160,1) and takes place in a stride of (1,5). As brittle cracks are more pronounced, they can be better detected using the method described. A larger kernel can therefore be used. As with "ductile" cracks, the shape of the crack is used for identification. Cracks are searched for along the horizontal. To reduce the noise, the Savitzky-Golay filter is applied again. The final identification of a crack is carried out analogously to the ductile cracks by the sequence of minima and maxima. To avoid errors, the identified cracks are checked and, if necessary, adjusted by the user as part of the evaluation process.

<figure><img src="image_781.png" /><figcaption>Figure 4 Processed images to identify brittle cracks based on image *i* (a) and total gradient image using image *i* and *i-2* (b)</figcaption></figure>

Finally, the sample is measured optically. The first image, i.e. the image in which there is no deformation, and the image that represents the time of crack initiation are used. The measurement is based on the distance between the shoulders of the sample. For this purpose, a Sobel filter is applied to both images to emphasise vertical edges. An average kernel of size (40,1) is subsequently applied to both filtered images along a line to take into account the shape of the shoulder as a vertical line and reduce noise. The edges of the shoulder can be identified via the maximum or minimum of the array starting from the beginning or end. Fig. 5 shows an example of this for the "brittle" crack. The edges of the sample shoulders are clearly recognisable in all images and were correctly identified by the algorithm. Finally, the strain is calculated using the measured sample lengths. The sample length itself is taken as a unitless proportion of the total image width. A critical strain which leads to failure of the oxide scale layer is calculated as the change in length ($l_1-l_0$) of its original length ($l_0$). As incorrect crack detection cannot be completely prevented, a final check is carried out by manual user control of the identified crack time and the determined line for measuring the change in length. For this purpose, the corresponding images are displayed after the evaluation. This was carried out for all evaluated tests, with the algorithm demonstrating a very high identification rate.
<figure><img src="image_774.png" /><figcaption>Figure 5 Optical measurement of length change for fracture initiation</figcaption></figure>

## Results

The tensile tests performed on defined oxidised samples were analysed using the evaluation tool. Samples showing an initial failure of the oxide layer were removed from the evaluation. The focus was on analysing the fracture behaviour and determining the characteristic critical strain at crack initiation. The Figure 6 shows the critical strain for different oxidation temperatures or test temperatures. In this case, the oxidation temperature was chosen to be the same as the test temperature. The strain rate was $0.01 \text{ s}^{-1}$.

<figure><img src="image_775.png" /><figcaption>Figure 6 Influence of temperature on critical strain at crack initiation for a strain rate of $0.01 \text{ s}^{-1}$</figcaption></figure>

The curve shows a decrease in the achievable critical strain of the oxide scale layer with increasing temperature. Low temperatures in the range of 900 °C - 950 °C indicate a ductile behaviour of the oxide scale layer. In contrast, high temperatures of above 1000 °C indicate brittle fracture behaviour. This behaviour can also be seen in the fracture images of the oxide scale layer. Fig. 7 shows a sample at the time of crack initiation (a) and at a later point in time (b) for a test temperature of 900 °C and a strain rate of $0.01 \text{ s}^{-1}$. It is clearly visible that several fine cracks form on the sample, but these do not increase significantly even with ongoing deformation. Brittle cracking or spallation of the oxide scale layer is also not observed. In addition, Fig. 7 shows a sample for a test temperature of 1100 °C for a strain rate of $0.01 \text{ s}^{-1}$ at the time of crack initiation (c) and at a later time (d). At the time of crack initiation, a single brittle crack can be seen running down the center of the entire sample. As the deformation progresses, a clear widening of the crack can be seen. Once a crack has formed, the oxide scale layer does not appear to undergo any further plastic deformation. This leads to an enlargement of the crack. Furthermore, this indicates a sliding of the oxide scale layer on the steel base substrate, which continues to deform plastically.
<figure><img src="image_782.png" /><figcaption>Figure 5 Optical measurement of length change for fracture initiation</figcaption></figure>

## Results

The tensile tests performed on defined oxidised samples were analysed using the evaluation tool. Samples showing an initial failure of the oxide layer were removed from the evaluation. The focus was on analysing the fracture behaviour and determining the characteristic critical strain at crack initiation. The Figure 6 shows the critical strain for different oxidation temperatures or test temperatures. In this case, the oxidation temperature was chosen to be the same as the test temperature. The strain rate was $0.01 \text{ s}^{-1}$.

<figure><img src="image_783.png" /><figcaption>Figure 6 Influence of temperature on critical strain at crack initiation for a strain rate of $0.01 \text{ s}^{-1}$</figcaption></figure>

The curve shows a decrease in the achievable critical strain of the oxide scale layer with increasing temperature. Low temperatures in the range of 900 °C - 950 °C indicate a ductile behaviour of the oxide scale layer. In contrast, high temperatures of above 1000 °C indicate brittle fracture behaviour. This behaviour can also be seen in the fracture images of the oxide scale layer. Fig. 7 shows a sample at the time of crack initiation (a) and at a later point in time (b) for a test temperature of 900 °C and a strain rate of $0.01 \text{ s}^{-1}$. It is clearly visible that several fine cracks form on the sample, but these do not increase significantly even with ongoing deformation. Brittle cracking or spallation of the oxide scale layer is also not observed. In addition, Fig. 7 shows a sample for a test temperature of 1100 °C for a strain rate of $0.01 \text{ s}^{-1}$ at the time of crack initiation (c) and at a later time (d). At the time of crack initiation, a single brittle crack can be seen running down the center of the entire sample. As the deformation progresses, a clear widening of the crack can be seen. Once a crack has formed, the oxide scale layer does not appear to undergo any further plastic deformation. This leads to an enlargement of the crack. Furthermore, this indicates a sliding of the oxide scale layer on the steel base substrate, which continues to deform plastically.
<figure><img src="image_776.png" /><figcaption>Figure 7 Ductile oxide scale failure at a test temperature of 900 °C at crack initiation (a) and at a later stage (b) as well as at a test temperature of 1100 °C at crack initiation (c) and at a later stage (d) for a strain rate of 0.01 s⁻¹</figcaption></figure>

Fig. 8 (a) shows the influence of the strain rate on the achievable strain within the oxide scale layer for a constant oxidation temperature of 900 °C. Even in this case, the test temperature was chosen to be the same as the oxidation temperature. It can be seen that for low strain rates in the range 0.01 s⁻¹ to 0.05 s⁻¹ no significant effect can be observed. However, when the strain rate is increased to 0.1 s⁻¹, there is a considerable decrease in the critical strain, which indicates a change in the fracture shape from ductile to brittle.

<figure><img src="image_777.png" /><figcaption>Figure 8 Influence of strain rate on oxide scale failure for a test temperature of 900 °C (a) and influence of test temperature on oxide scale failure for an oxidation temperature of 1050 °C and a strain rate of 0.01 s⁻¹ (b)</figcaption></figure>

Finally, the influence of different test temperatures with constant oxidation parameters was investigated. This case illustrates oxidation of a heated semi-finished product, for example during transfer and cooling of the oxide scale by air or contact with the cold tools. Furthermore, in this case the constant oxidation conditions result in an oxide scale layer with the same oxide scale layer thickness and composition. For this purpose, an oxidation temperature of 1050 °C was selected and the samples were cooled to test temperatures between 900 °C and 450 °C in accordance with the procedure described. All tests were carried out at a strain rate of 0.01 s⁻¹. Compared to the other samples, these tests therefore always show a comparable oxide scale layer thickness and composition. The influence of different test temperatures on the achievable critical strain is shown in Fig. 9 (b). In the test temperature range between 600 °C and 950 °C, no significant influence of the temperature on the critical strain can be observed. However, a temperature of 400 °C leads to a significant decrease in the achievable strain. Overall, the critical degrees of deformation are well below the values observed for higher temperatures (Fig. 6). The degrees of deformation achieved
<figure><img src="image_784.png" /><figcaption>Figure 7 Ductile oxide scale failure at a test temperature of 900 °C at crack initiation (a) and at a later stage (b) as well as at a test temperature of 1100 °C at crack initiation (c) and at a later stage (d) for a strain rate of 0.01 s⁻¹</figcaption></figure>

Fig. 8 (a) shows the influence of the strain rate on the achievable strain within the oxide scale layer for a constant oxidation temperature of 900 °C. Even in this case, the test temperature was chosen to be the same as the oxidation temperature. It can be seen that for low strain rates in the range 0.01 s⁻¹ to 0.05 s⁻¹ no significant effect can be observed. However, when the strain rate is increased to 0.1 s⁻¹, there is a considerable decrease in the critical strain, which indicates a change in the fracture shape from ductile to brittle.

<figure><img src="image_785.png" /><figcaption>Figure 8 Influence of strain rate on oxide scale failure for a test temperature of 900 °C (a) and influence of test temperature on oxide scale failure for an oxidation temperature of 1050 °C and a strain rate of 0.01 s⁻¹ (b)</figcaption></figure>

Finally, the influence of different test temperatures with constant oxidation parameters was investigated. This case illustrates oxidation of a heated semi-finished product, for example during transfer and cooling of the oxide scale by air or contact with the cold tools. Furthermore, in this case the constant oxidation conditions result in an oxide scale layer with the same oxide scale layer thickness and composition. For this purpose, an oxidation temperature of 1050 °C was selected and the samples were cooled to test temperatures between 900 °C and 450 °C in accordance with the procedure described. All tests were carried out at a strain rate of 0.01 s⁻¹. Compared to the other samples, these tests therefore always show a comparable oxide scale layer thickness and composition. The influence of different test temperatures on the achievable critical strain is shown in Fig. 9 (b). In the test temperature range between 600 °C and 950 °C, no significant influence of the temperature on the critical strain can be observed. However, a temperature of 400 °C leads to a significant decrease in the achievable strain. Overall, the critical degrees of deformation are well below the values observed for higher temperatures (Fig. 6). The degrees of deformation achieved
indicate a brittle failure of the oxide scale. This is also illustrated by the images of the sample at 900 °C (a, b) and 480 °C (c, d) shown in Fig. 10.

<figure><img src="image_778.png" /><figcaption>Figure 9 Oxide failure at a test temperature of 900 °C at crack initiation (a) and at a later stage (b) and at a test temperature of 450 °C at crack initiation (c) and at a later stage (d) for an oxidation temperature of 1050 °C and a atrain rate of 0.01 s-1</figcaption></figure>

Both test temperatures show a brittle crack pattern and an enlargement of the crack with increasing mechanical load. At a test temperature of 450 °C, several fine cracks can be seen in the sample, which increase in size continuously. This indicates a lower strength of the oxide scale or a stronger bonding of the oxide scale to the steel substrate. Unlike at 900 °C, it is not possible for a larger oxide scale layer to slide off.

## Summary

In this work, the damage behaviour of oxide scale on C45 steel was analysed under hot forging-relevant process conditions. For the local and temporally resolved investigation of the damage, flat tensile tests were carried out on specifically oxidized samples, which were recorded with an optical measuring system. An evaluation algorithm based on computer vision approaches was developed to evaluate the recorded image series. The subsequent analysis of the tests carried out showed that the oxide scale layer exhibits different forms of damage depending on the temperature and strain rate. Low oxidation temperatures result in a rather ductile behaviour of the oxide scale layer, whereas higher oxidation temperatures tend to lead to a brittle damage behaviour. A more ductile behaviour can also be observed at low strain rates. Higher strain rates result in a brittle fracture of the oxide scale layer. However, the work of [6] reveals that even with the same oxidation time, different oxidation temperatures lead to different oxide scale layer thicknesses and compositions. Thus, in further work, the oxide scale layer thickness and composition are therefore to be calculated based on the work of [16] and correlated with the results obtained on damage behaviour. Tests with comparable oxidation parameters and thus comparable oxide scale layer thicknesses and compositions also showed an influence of the test temperature. A decrease of achievable critical strain was observed for reduced test temperatures. In order to make more holistic statements about the effects and to analyse interactions between individual effects in more detail, the test matrix is to be expanded in further work.

## Acknowledgement

The authors thank the German Research Foundation (DFG) for the financial support of the project "General modelling of material behaviour and surface modifications for FEM analysis of the die forging of carbon steels" (Project number 316273316).
indicate a brittle failure of the oxide scale. This is also illustrated by the images of the sample at 900 °C (a, b) and 480 °C (c, d) shown in Fig. 10.

<figure><img src="image_786.png" /><figcaption>Figure 9 Oxide failure at a test temperature of 900 °C at crack initiation (a) and at a later stage (b) and at a test temperature of 450 °C at crack initiation (c) and at a later stage (d) for an oxidation temperature of 1050 °C and a atrain rate of 0.01 s-1</figcaption></figure>

Both test temperatures show a brittle crack pattern and an enlargement of the crack with increasing mechanical load. At a test temperature of 450 °C, several fine cracks can be seen in the sample, which increase in size continuously. This indicates a lower strength of the oxide scale or a stronger bonding of the oxide scale to the steel substrate. Unlike at 900 °C, it is not possible for a larger oxide scale layer to slide off.

## Summary

In this work, the damage behaviour of oxide scale on C45 steel was analysed under hot forging-relevant process conditions. For the local and temporally resolved investigation of the damage, flat tensile tests were carried out on specifically oxidized samples, which were recorded with an optical measuring system. An evaluation algorithm based on computer vision approaches was developed to evaluate the recorded image series. The subsequent analysis of the tests carried out showed that the oxide scale layer exhibits different forms of damage depending on the temperature and strain rate. Low oxidation temperatures result in a rather ductile behaviour of the oxide scale layer, whereas higher oxidation temperatures tend to lead to a brittle damage behaviour. A more ductile behaviour can also be observed at low strain rates. Higher strain rates result in a brittle fracture of the oxide scale layer. However, the work of [6] reveals that even with the same oxidation time, different oxidation temperatures lead to different oxide scale layer thicknesses and compositions. Thus, in further work, the oxide scale layer thickness and composition are therefore to be calculated based on the work of [16] and correlated with the results obtained on damage behaviour. Tests with comparable oxidation parameters and thus comparable oxide scale layer thicknesses and compositions also showed an influence of the test temperature. A decrease of achievable critical strain was observed for reduced test temperatures. In order to make more holistic statements about the effects and to analyse interactions between individual effects in more detail, the test matrix is to be expanded in further work.

## Acknowledgement

The authors thank the German Research Foundation (DFG) for the financial support of the project "General modelling of material behaviour and surface modifications for FEM analysis of the die forging of carbon steels" (Project number 316273316).
## References

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[6] G. Korpała, M. Ullmann, M. Graf, H. Wester, A. Bouguecha, B. Awiszus, B.-A. Behrens, and R. Kawalla, Modelling the influence of carbon content on material behavior during forging. DOI=10.1063/1.5008226.

[7] Graf, M. (2013. Modellierung des Umformverhaltens von Zunder entlang der Prozesskette Warmband. Freiberger Forschungshefte. B, Werkstofftechnologie 353. Technische Universität Bergakademie Freiberg, Freiberg.

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[10] L. Suárez, Y. Houbaert, X. V. Eynde, and R. Colás, High temperature deformation of oxide scale, Corrosion Science 51 (2) (2009) 309-315. https://doi.org/10.1016/j.corsci.2008.10.027

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## References

[1] B.-A. Behrens, R. Kawalla, B. Awiszus, A. Bouguecha, M. Ullmann, M. Graf, C. Bonk, A. Chugreev, and H. Wester, Numerical Investigation of the Oxide Scale Deformation Behaviour with Consideration of Carbon Content during Hot Forging, Procedia Engineering 207 (2017) 526-531. https://doi.org/10.1016/j.proeng.2017.10.816

[2] R. Kawalla and F. Steinert, Untersuchung des Einflusses von Prozessparametern in der Fertigstraße auf die Tertiärzunderausbildung, Materialwissenschaft und Werkstofftechnik 1 (38) (2007) 36-42. https://doi.org/10.1002/mawe.200600075

[3] L. Luong and T. Heijkoop, The influence of scale on friction in hot metal working, Wear 71 (1) (1981) 93-102. https://doi.org/10.1016/0043-1648(81)90142-3

[4] M. Schütze, Mechanical properties of oxide scales, Oxid Met 44 (1-2) (1995) 29-61. https://doi.org/10.1007/BF01046722

[5] H. Utsunomiya, S. Doi, K. Hara, T. Sakai, and S. Yanagi, Deformation of oxide scale on steel surface during hot rolling, CIRP Annals 58 (1) (2009) 271-274. https://doi.org/10.1016/j.cirp.2009.03.050

[6] G. Korpała, M. Ullmann, M. Graf, H. Wester, A. Bouguecha, B. Awiszus, B.-A. Behrens, and R. Kawalla, Modelling the influence of carbon content on material behavior during forging. DOI=10.1063/1.5008226.

[7] Graf, M. (2013. Modellierung des Umformverhaltens von Zunder entlang der Prozesskette Warmband. Freiberger Forschungshefte. B, Werkstofftechnologie 353. Technische Universität Bergakademie Freiberg, Freiberg.

[8] M. Graf and R. Kawalla, Scale Behaviour and Deformation Properties of Oxide Scale during Hot Rolling of Steel, KEM 504-506 (2012) 199-204. https://doi.org/10.4028/www.scientific.net/KEM.504-506.199

[9] J. Favergeon, G. Moulin, A. Makni, and L. Lahoche, The Effect of Oxide Scale on the Mechanical Behavior of Low Alloyed Steel at High Temperature, MSF 522-523 (2006) 401-408. https://doi.org/10.4028/www.scientific.net/MSF.522-523.401

[10] L. Suárez, Y. Houbaert, X. V. Eynde, and R. Colás, High temperature deformation of oxide scale, Corrosion Science 51 (2) (2009) 309-315. https://doi.org/10.1016/j.corsci.2008.10.027

[11] M. Schütze, P. F. Tortorelli, and I. G. Wright, Development of a Comprehensive Oxide Scale Failure Diagram, Oxid Met 73 (3-4) (2010) 389-418. https://doi.org/10.1007/s11085-009-9185-7

[12] M. Krzyzanowski and J. H. Beynon, The tensile failure of mild steel oxides under hot rolling conditions, Steel Research 70 (1) (1999) 22-27. https://doi.org/10.1002/srin.199905596

[13] W. Gao, X. Zhang, L. Yang, and H. Liu (2010, An improved Sobel edge detection. In 3rd International Conference on Computer Science and Information Technology. IEEE) 67-71 (2010). https://doi.org/10.1109/ICCSIT.2010.5563693

[14] A. Savitzky and M. J. E. Golay, Smoothing and Differentiation of Data by Simplified Least Squares Procedures, Anal. Chem. 36 (8) (1964) 1627-1639. https://doi.org/10.1021/ac60214a047

[15] R. Jardim and F. Morgado-Dias, Savitzky-Golay filtering as image noise reduction with sharp color reset, Microprocessors and Microsystems 74 (2020). https://doi.org/10.1016/j.micpro.2020.103006

[16] T. Bergelt, M. Graf, J. N. Hunze, B.-A. Behrens, and T. Lampke, Investigation of scale properties and layer growth depending on the carbon and chromium content in steel, European Oxide Scale Conference (2022). https://doi.org/10.5281/zenodo.10138002
# Numerical process design for the production of a hybrid die made of tool steel X38CrMoV5.3 and inconel 718

SIRING Janina¹,ᵃ*, HEINE Christoph¹, TILL Michael¹, WESTER Hendrik¹, UHE Johanna¹, BEHRENS Bernd-Arno¹ and BRUNOTTE Kai¹,ᵇ

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

ᵃsiring@ifum.uni-hannover.de, ᵇbrunotte@ifum.uni-hannover.de

**Keywords:** Tailored Forming, Hybrid Forging Dies, X38CrMoV5.3, Inconel 718, Numerical Die Design

**Abstract.** Dies used in hot forging are subjected to high cyclic thermo-mechanical loads, which lead to die failure. There are various options for increasing the service life of these dies, for example coatings or heat treatments. Another possibility is to adapt the choice of material, which is the focus of this work. For example, the nickel-based alloy Inconel has a higher strength at elevated temperatures compared to tool steel. However, Inconel is difficult to manufacture and has higher material costs. For this reason, a new process design for the production of a hybrid die consisting of Inconel 718 and tool steel X38CrMoV5.3 is presented within this work. To produce the hybrid dies, the two materials are first friction welded and then formed using hot forging. In addition to the numerical process design, experimental tests are also carried out to manufacture such hybrid dies. Furthermore, a numerical parameter study is done to determine the influence of the forging temperature, the forging speed and the initial Inconel thickness on the process parameters. It can be shown that the production of hybrid dies is possible by using the Tailored Forming process chain. The influencing factors investigated change the required press force and also the material distribution of the Inconel in the hybrid die produced. In the future, further experimental tests will be carried out to determine the service life of the hybrid dies.

**Introduction**

In hot forging, the dies are exposed to high thermal and mechanical cyclic loads. These can lead to failure, particular in the die surfaces [1]. Die failure causes high costs due to the replacement as well as production downtimes [2]. The total resulting costs depend not only on the failure of the die, but also on the logistics costs. Therefore, the aim should be to increase the service life, but not to generate higher logistics planning through complicated die production [3]. Various approaches exist to increase the die service life. On the one hand, the loads that occur can be reduced by improving the tribological conditions or adapting the die geometry [2]. On the other hand, the load resistance can be increased by surface modification [4] or choice of material [5].

A material with a high potential to increase the load resistance of forging dies is Inconel. Compared to tool steel, Inconel offers high mechanical properties, especially at elevated load temperatures. However, Inconel is more expensive than tool steel and difficult to manufacture [6]. Therefore, within this work a production process concept for the manufacturing of a hybrid forging die is presented and investigated. The hybrid die consists of the conventional tool steel X38CrMoV5.3 (AISI H10 / DIN: 1.2367) and the Inconel alloy 718 (EN NiCr19NbMo / 2.4668). Inconel is used for the near surface regions to increase the die life due to its high strength at elevated temperatures. The significantly less expensive tool steel is used in less critical areas as base material. To counteract the low machinability of Inconel, the hybrid dies are produced by means of die forging. According to the Tailored Forming concept [7], friction welding is used to join both materials before the forging process.

<figure><img src="image_779.png" /></figure>

Content from this work may be used under the terms of the Creative Commons Attribution 3.0 license. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. Published under license by Materials Research Forum LLC.
# Numerical process design for the production of a hybrid die made of tool steel X38CrMoV5.3 and inconel 718

SIRING Janina¹,ᵃ*, HEINE Christoph¹, TILL Michael¹, WESTER Hendrik¹, UHE Johanna¹, BEHRENS Bernd-Arno¹ and BRUNOTTE Kai¹,ᵇ

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

ᵃsiring@ifum.uni-hannover.de, ᵇbrunotte@ifum.uni-hannover.de

**Keywords:** Tailored Forming, Hybrid Forging Dies, X38CrMoV5.3, Inconel 718, Numerical Die Design

**Abstract.** Dies used in hot forging are subjected to high cyclic thermo-mechanical loads, which lead to die failure. There are various options for increasing the service life of these dies, for example coatings or heat treatments. Another possibility is to adapt the choice of material, which is the focus of this work. For example, the nickel-based alloy Inconel has a higher strength at elevated temperatures compared to tool steel. However, Inconel is difficult to manufacture and has higher material costs. For this reason, a new process design for the production of a hybrid die consisting of Inconel 718 and tool steel X38CrMoV5.3 is presented within this work. To produce the hybrid dies, the two materials are first friction welded and then formed using hot forging. In addition to the numerical process design, experimental tests are also carried out to manufacture such hybrid dies. Furthermore, a numerical parameter study is done to determine the influence of the forging temperature, the forging speed and the initial Inconel thickness on the process parameters. It can be shown that the production of hybrid dies is possible by using the Tailored Forming process chain. The influencing factors investigated change the required press force and also the material distribution of the Inconel in the hybrid die produced. In the future, further experimental tests will be carried out to determine the service life of the hybrid dies.

**Introduction**

In hot forging, the dies are exposed to high thermal and mechanical cyclic loads. These can lead to failure, particular in the die surfaces [1]. Die failure causes high costs due to the replacement as well as production downtimes [2]. The total resulting costs depend not only on the failure of the die, but also on the logistics costs. Therefore, the aim should be to increase the service life, but not to generate higher logistics planning through complicated die production [3]. Various approaches exist to increase the die service life. On the one hand, the loads that occur can be reduced by improving the tribological conditions or adapting the die geometry [2]. On the other hand, the load resistance can be increased by surface modification [4] or choice of material [5].

A material with a high potential to increase the load resistance of forging dies is Inconel. Compared to tool steel, Inconel offers high mechanical properties, especially at elevated load temperatures. However, Inconel is more expensive than tool steel and difficult to manufacture [6]. Therefore, within this work a production process concept for the manufacturing of a hybrid forging die is presented and investigated. The hybrid die consists of the conventional tool steel X38CrMoV5.3 (AISI H10 / DIN: 1.2367) and the Inconel alloy 718 (EN NiCr19NbMo / 2.4668). Inconel is used for the near surface regions to increase the die life due to its high strength at elevated temperatures. The significantly less expensive tool steel is used in less critical areas as base material. To counteract the low machinability of Inconel, the hybrid dies are produced by means of die forging. According to the Tailored Forming concept [7], friction welding is used to join both materials before the forging process.

<figure><img src="image_787.png" /></figure>

Content from this work may be used under the terms of the Creative Commons Attribution 3.0 license. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. Published under license by Materials Research Forum LLC.
Numerical simulation is used to design the manufacturing process for the hybrid dies. Due to the different material properties and the difficulty of reworking Inconel, it is essential to design the process close to the final geometry. In addition, it must be ensured that a sufficient Inconel layer is present locally in accordance with the expected loads. The usability of finite element (FE) simulation has already been demonstrated in the work of Bouguecha et al. [2] for the design of forging dies and Renhof [5] for the prediction of the material behaviour as well as for the dimensioning of hybrid forging processes by Behrens et al. [7].

The advantage of a hybrid material combination has been investigated by Landgrebe et al. [8]. Here, an Inconel layer was used to reduce the thermal and corrosive load as well as to improve frictional resistance on the component. The hybrid material combination was produced by selective laser melting (SLM) process. The potential of the process chain, consisting of joining the two materials and then combined forging, has already been partially investigated. For example, the process chain for the production of hybrid shafts [7] or bearing bushings [9] was analysed. In these investigations, the focus was on the material combination of steel and aluminium as well as on the production of hybrid components. The production of dies for application in hot forging processes using a forming process was investigated in the work of Behrens et al. [10]. However, only the use of mono-material was considered. The forming process results in a higher surface hardness compared to a conventionally manufactured die.

The general feasibility of the combined forging of pre-joined semi-finished workpiece made of tool steel and Inconel has already been demonstrated in initial numerical and experimental investigations for a simplified geometry [11]. However, the production of a specific die shape has not yet been analysed.

This paper focuses on the numerical design and analysis of a production process for a hybrid die using a mandrel. In particular, the material flow of the Inconel layer must be analysed in order to ensure a suitable distribution of Inconel in the final component. In this context, the differing deformation behaviour of the used tool steel and Inconel as well as the die geometry influences the final material distribution of the hybrid die. Furthermore, the higher strength of Inconel even at elevated temperatures leads to higher die loads, which also needs to be taken into account during process design. The numerical model was validated by experimental forging tests and subsequently used for a numerical parameter study. Within the numerical study, the influence of the initial Inconel layer thickness as well as process parameters like press speed or temperature on the final material distribution and the required process force were analysed. Based on the results a possible process parameter setup was defined. The focus was on realising the lowest possible Inconel thickness in order to keep the material costs as low as possible.

## Procedure

This work focuses on the process design and analysis for the production of a hybrid mandrel die geometry. The hybrid dies are to be made of tool steel X38CrMoV5.3 and Inconel 718, whereby the Inconel is to cover the zone most affected by the temperature and thus reduce the loads on the softer tool steel. For an efficient process design FE simulation using Simufact Forming v16.0 is used.

**Definition of required Inconel thickness.** In order to determine the required Inconel thickness, an FE simulation of the forging process was carried out with the chosen die geometry consisting of the mono-material X38CrMoV5.3. An overview over the used numerical process model as well as the considered process conditions can be found in [12]. The material model used was taken from the work of [11]. Based on the calculated local temperature distribution in the near surface regions the Inconel layer thickness is defined.
Numerical simulation is used to design the manufacturing process for the hybrid dies. Due to the different material properties and the difficulty of reworking Inconel, it is essential to design the process close to the final geometry. In addition, it must be ensured that a sufficient Inconel layer is present locally in accordance with the expected loads. The usability of finite element (FE) simulation has already been demonstrated in the work of Bouguecha et al. [2] for the design of forging dies and Renhof [5] for the prediction of the material behaviour as well as for the dimensioning of hybrid forging processes by Behrens et al. [7].

The advantage of a hybrid material combination has been investigated by Landgrebe et al. [8]. Here, an Inconel layer was used to reduce the thermal and corrosive load as well as to improve frictional resistance on the component. The hybrid material combination was produced by selective laser melting (SLM) process. The potential of the process chain, consisting of joining the two materials and then combined forging, has already been partially investigated. For example, the process chain for the production of hybrid shafts [7] or bearing bushings [9] was analysed. In these investigations, the focus was on the material combination of steel and aluminium as well as on the production of hybrid components. The production of dies for application in hot forging processes using a forming process was investigated in the work of Behrens et al. [10]. However, only the use of mono-material was considered. The forming process results in a higher surface hardness compared to a conventionally manufactured die.

The general feasibility of the combined forging of pre-joined semi-finished workpiece made of tool steel and Inconel has already been demonstrated in initial numerical and experimental investigations for a simplified geometry [11]. However, the production of a specific die shape has not yet been analysed.

This paper focuses on the numerical design and analysis of a production process for a hybrid die using a mandrel. In particular, the material flow of the Inconel layer must be analysed in order to ensure a suitable distribution of Inconel in the final component. In this context, the differing deformation behaviour of the used tool steel and Inconel as well as the die geometry influences the final material distribution of the hybrid die. Furthermore, the higher strength of Inconel even at elevated temperatures leads to higher die loads, which also needs to be taken into account during process design. The numerical model was validated by experimental forging tests and subsequently used for a numerical parameter study. Within the numerical study, the influence of the initial Inconel layer thickness as well as process parameters like press speed or temperature on the final material distribution and the required process force were analysed. Based on the results a possible process parameter setup was defined. The focus was on realising the lowest possible Inconel thickness in order to keep the material costs as low as possible.

## Procedure

This work focuses on the process design and analysis for the production of a hybrid mandrel die geometry. The hybrid dies are to be made of tool steel X38CrMoV5.3 and Inconel 718, whereby the Inconel is to cover the zone most affected by the temperature and thus reduce the loads on the softer tool steel. For an efficient process design FE simulation using Simufact Forming v16.0 is used.

**Definition of required Inconel thickness.** In order to determine the required Inconel thickness, an FE simulation of the forging process was carried out with the chosen die geometry consisting of the mono-material X38CrMoV5.3. An overview over the used numerical process model as well as the considered process conditions can be found in [12]. The material model used was taken from the work of [11]. Based on the calculated local temperature distribution in the near surface regions the Inconel layer thickness is defined.
**Numerical setup and parameter study.** Fig. 1b shows a schematic representation of the simulation model of the forging process for manufacturing the hybrid die. The punch, ejector, reinforcement ring and shoulder insert consist of X37CrMoV5.1. The dies are not heated during the forging process, therefore an initial die temperature of 20 °C was assumed for the simulation. Ejector and punch are modelled as rigid bodies with heat conduction. The highly stressed sections of the mould, the shoulder insert and the reinforcement ring were modelled as deformable parts. The hybrid semi-finished workpiece consisting of the tool steel X38CrMoV5.3 and the Inconel 718 was simulated with a tie contact in order to depict the material bonded connection produced by means of friction welding. Material data like temperature and strain rate dependent flow curves for both materials were experimental determined in previous work of the authors [11]. Before the simulation of the forging process, the hybrid semi-finished product was heated up and a subsequent cooling simulation with a duration of 20 s was carried out. It was heated up to the temperature corresponding to the subsequent furnace temperature. The individual simulations are used to reproduce the experimental process conditions, whereby a cooling time equal to the transfer time from the furnace to the press was selected. In all simulations (explicit time integration), the dies were meshed with "Quadtree" elements and the semi-finished product with "Advancing Front Quad" elements.

<figure><img src="image_780.png" /><figcaption>Fig. 1. (a) Schematic sequence of the process chain for producing a hybrid die and (b) Simulation model.</figcaption></figure>

A combined friction model of Coulomb and Tresca was selected with a friction coefficient $\mu$ of 0.15 and a friction factor $m$ of 0.3, which was defined in accordance with [2]. With these friction parameters, very good agreement between simulation and experiment was achieved in this project as well as in other projects. The heat transfer coefficient between Inconel and tool steel was set to 5 W/(m²K) based on data from literature [13]. In order to take into account the different changes in length of the two materials both during the heating of the hybrid semi-finished workpiece and the subsequent hot forging the coefficient of thermal expansion was determined on the basis of literature values. For Inconel, a value averaged between the work of Agazhanov et al. [14] and Dye et al. [15] was defined. For tool steel X38CrMoV5.3 the value of [16] was used. Punch movement is modelled using a hydraulic press with a constant press speed, which will also be used for the experimental forging tests. As part of the full factorial parameter study, the press speed (5 / 10 / 15 mm/s) and the Inconel thickness (3 / 5 / 7 mm) were varied. The height of the tool steel is adjusted so that the overall height of the hybrid semi-finished workpiece remains constant at
**Numerical setup and parameter study.** Fig. 1b shows a schematic representation of the simulation model of the forging process for manufacturing the hybrid die. The punch, ejector, reinforcement ring and shoulder insert consist of X37CrMoV5.1. The dies are not heated during the forging process, therefore an initial die temperature of 20 °C was assumed for the simulation. Ejector and punch are modelled as rigid bodies with heat conduction. The highly stressed sections of the mould, the shoulder insert and the reinforcement ring were modelled as deformable parts. The hybrid semi-finished workpiece consisting of the tool steel X38CrMoV5.3 and the Inconel 718 was simulated with a tie contact in order to depict the material bonded connection produced by means of friction welding. Material data like temperature and strain rate dependent flow curves for both materials were experimental determined in previous work of the authors [11]. Before the simulation of the forging process, the hybrid semi-finished product was heated up and a subsequent cooling simulation with a duration of 20 s was carried out. It was heated up to the temperature corresponding to the subsequent furnace temperature. The individual simulations are used to reproduce the experimental process conditions, whereby a cooling time equal to the transfer time from the furnace to the press was selected. In all simulations (explicit time integration), the dies were meshed with "Quadtree" elements and the semi-finished product with "Advancing Front Quad" elements.

<figure><img src="image_788.png" /><figcaption>Fig. 1. (a) Schematic sequence of the process chain for producing a hybrid die and (b) Simulation model.</figcaption></figure>

A combined friction model of Coulomb and Tresca was selected with a friction coefficient $\mu$ of 0.15 and a friction factor $m$ of 0.3, which was defined in accordance with [2]. With these friction parameters, very good agreement between simulation and experiment was achieved in this project as well as in other projects. The heat transfer coefficient between Inconel and tool steel was set to 5 W/(m$^2$K) based on data from literature [13]. In order to take into account the different changes in length of the two materials both during the heating of the hybrid semi-finished workpiece and the subsequent hot forging the coefficient of thermal expansion was determined on the basis of literature values. For Inconel, a value averaged between the work of Agazhanov et al. [14] and Dye et al. [15] was defined. For tool steel X38CrMoV5.3 the value of [16] was used. Punch movement is modelled using a hydraulic press with a constant press speed, which will also be used for the experimental forging tests. As part of the full factorial parameter study, the press speed (5 / 10 / 15 mm/s) and the Inconel thickness (3 / 5 / 7 mm) were varied. The height of the tool steel is adjusted so that the overall height of the hybrid semi-finished workpiece remains constant at
33 mm. Furthermore, the selected furnace temperature (900 / 1,000 / 1,100 °C) was varied as part of the parameter study.

**Experimental setup.** Based on the numerical results experimental forging tests for the production of a hybrid die was carried out using friction-welded hybrid semi-finished workpieces. A schematic process chain for the manufacture of a hybrid die is shown in Fig. 1a. The pre-joined hybrid semi-finished workpiece consist of an Inconel layer with an initial thickness of 5 mm whereby the initial height of the tool steel was 28 mm. The hybrid semi-finished workpiece was preheated to a temperature of 1,000 °C within a furnace. Heating time was set to 15 min with a 5 min hold time to ensure homogenous temperature distribution. The transfer time between the furnace and the press was measured to 20 s. Tests were carried out on a hydraulic press and a constant press speed of 10 mm/s was set. After successful forging of the hybrid semi-finished workpiece, the hybrid die was measured optically. The optical measurement was converted into a digital model of the produced hybrid die and compared to the calculated geometry. For later use as a die in hot forging, the surface still needs to be reworked. The post-processing is based on the work of [17].

**Results**

The results are differentiated between the verification of the simulation results based on experimental data and the parameter study. Firstly, the comparison between simulation and experiment is discussed in more detail. After successful verification of the simulation model, the results from the parameter study are presented.

**Minimum Inconel layer thickness.** The numerical calculated temperature distribution shown in Fig. 2a indicates that the highest thermal loads occur at the mandrel. With increasing radius to the die centre, the temperatures near the surface decrease. In the centre of the mandrel, an increased temperature can be observed to a depth of 2.4 mm. Near the mandrel radius a temperature influence zone of up to 2.2 mm from the surface was determined. However, these region shows the highest temperature of up to 573 °C. Based on these findings, the final Inconel layer should have a thickness of at least 2.4 mm in the area of the highest thermal load after hybrid forging. At the same time, an Inconel layer is not required everywhere on the die surface. As a result, a separated die concept was developed (Fig. 2b). The hybrid die will consist of a conventional outer section and an inner hybrid section, whereby the inner hybrid mandrel is fixed (shrink fit and press-fit connection) in the conventional section. This will simplify the production of the hybrid inserts and any possible removal of the highly loaded mandrel can take place separately in the subsequent use. In addition, a material-efficient realisation can be carried out in the die.

<figure><img src="image_781.png" /><figcaption>Fig. 2. Result of the temperature analysis of the mono-material die (a) and schematic illustration of the proposed hybrid die concept (b).</figcaption></figure>
33 mm. Furthermore, the selected furnace temperature (900 / 1,000 / 1,100 °C) was varied as part of the parameter study.

**Experimental setup.** Based on the numerical results experimental forging tests for the production of a hybrid die was carried out using friction-welded hybrid semi-finished workpieces. A schematic process chain for the manufacture of a hybrid die is shown in Fig. 1a. The pre-joined hybrid semi-finished workpiece consist of an Inconel layer with an initial thickness of 5 mm whereby the initial height of the tool steel was 28 mm. The hybrid semi-finished workpiece was preheated to a temperature of 1,000 °C within a furnace. Heating time was set to 15 min with a 5 min hold time to ensure homogenous temperature distribution. The transfer time between the furnace and the press was measured to 20 s. Tests were carried out on a hydraulic press and a constant press speed of 10 mm/s was set. After successful forging of the hybrid semi-finished workpiece, the hybrid die was measured optically. The optical measurement was converted into a digital model of the produced hybrid die and compared to the calculated geometry. For later use as a die in hot forging, the surface still needs to be reworked. The post-processing is based on the work of [17].

## Results

The results are differentiated between the verification of the simulation results based on experimental data and the parameter study. Firstly, the comparison between simulation and experiment is discussed in more detail. After successful verification of the simulation model, the results from the parameter study are presented.

**Minimum Inconel layer thickness.** The numerical calculated temperature distribution shown in Fig. 2a indicates that the highest thermal loads occur at the mandrel. With increasing radius to the die centre, the temperatures near the surface decrease. In the centre of the mandrel, an increased temperature can be observed to a depth of 2.4 mm. Near the mandrel radius a temperature influence zone of up to 2.2 mm from the surface was determined. However, these region shows the highest temperature of up to 573 °C. Based on these findings, the final Inconel layer should have a thickness of at least 2.4 mm in the area of the highest thermal load after hybrid forging. At the same time, an Inconel layer is not required everywhere on the die surface. As a result, a separated die concept was developed (Fig. 2b). The hybrid die will consist of a conventional outer section and an inner hybrid section, whereby the inner hybrid mandrel is fixed (shrink fit and press-fit connection) in the conventional section. This will simplify the production of the hybrid inserts and any possible removal of the highly loaded mandrel can take place separately in the subsequent use. In addition, a material-efficient realisation can be carried out in the die.

<figure><img src="image_789.png" /><figcaption>Fig. 2. Result of the temperature analysis of the mono-material die (a) and schematic illustration of the proposed hybrid die concept (b).</figcaption></figure>
**Numerical process design and validation.** Using the described FE model, the process was numerical depicted. Initially, an Inconel layer thickness of 5 mm was considered. The temperature of the hybrid semi-finished workpiece was set to 1,000 °C and a press speed of 10 mm/s was assumed. The result of the simulation at 100 % forging with the punch released is shown in Fig. 3a. A complete filling of the die can be recognised. The Inconel shows the highest value of plastic strain of the hybrid semi-finished workpiece located in the upper area, where it is deformed by the shoulder insert. This plastic strain is evenly distributed in the tool steel at the point mentioned above. Within the tool steel, the plastic strain is rather low.

The load on the dies during the forging process is shown in Fig. 3b. It can be seen that the shoulder insert is subjected to the highest load. The maximum value of the equivalent stress v. Mises is almost 3,000 MPa. These values can be observed primarily in the upper radius, where the highest value of plastic strain of Inconel is located, and in the lower radius of the shoulder insert. The reinforcement ring shows equivalent stress values of 2,000-2,100 MPa. These are in the area where the hybrid semi-finished workpiece has contact and is upset. The numerical results show that the production of hybrid dies is possible using this set of parameters. These determined stress values are above the mechanical properties like yield strength of the material selected for the component [18]. This local exceeding of the yield strength can lead to plastic deformation of the die. As the hot forging process is intended to produce dies for later use in hot forging and as little material as possible should be used, it is necessary to produce the mandrels close to the final contour during forging. To ensure this, it is necessary to estimate the probability of plastic deformation in the die. For this reason, the equivalent stresses are analysed in this context. Due to the arrangement of a reinforcement ring around the shoulder insert, the material can be prevented from plastic deformation. Nevertheless, the aim of this process should be to minimise the load on the die sections.

<figure><img src="image_782.png" /><figcaption>Fig. 3. Result of the numerical process design (a) and the die load at 5 mm Inconel thickness, 10 mm/s press speed and 1,000 °C furnace temperature (b).</figcaption></figure>

Subsequently, experimental forging tests were performed. During the tests a maximum press force of 1,960-2,000 kN was reached. A nearly similar maximum force of 2,050 kN was calculated by numerical simulation. Overall, the experimental and numerical force-displacement curves show good agreement (Fig. 4a). The characteristic ranges and the maximum force are correctly depicted in the simulation. Deviations in the force curve at high stroke values can be attributed to elastic deflection of the press system in reality. Further, a geometric verification was carried out. For this purpose, the experimentally produced die was optically measured and a surface file was exported. A comparison between experimental measured and numerical calculated die geometry is shown in
**Numerical process design and validation.** Using the described FE model, the process was numerical depicted. Initially, an Inconel layer thickness of 5 mm was considered. The temperature of the hybrid semi-finished workpiece was set to 1,000 °C and a press speed of 10 mm/s was assumed. The result of the simulation at 100 % forging with the punch released is shown in Fig. 3a. A complete filling of the die can be recognised. The Inconel shows the highest value of plastic strain of the hybrid semi-finished workpiece located in the upper area, where it is deformed by the shoulder insert. This plastic strain is evenly distributed in the tool steel at the point mentioned above. Within the tool steel, the plastic strain is rather low.

The load on the dies during the forging process is shown in Fig. 3b. It can be seen that the shoulder insert is subjected to the highest load. The maximum value of the equivalent stress v. Mises is almost 3,000 MPa. These values can be observed primarily in the upper radius, where the highest value of plastic strain of Inconel is located, and in the lower radius of the shoulder insert. The reinforcement ring shows equivalent stress values of 2,000-2,100 MPa. These are in the area where the hybrid semi-finished workpiece has contact and is upset. The numerical results show that the production of hybrid dies is possible using this set of parameters. These determined stress values are above the mechanical properties like yield strength of the material selected for the component [18]. This local exceeding of the yield strength can lead to plastic deformation of the die. As the hot forging process is intended to produce dies for later use in hot forging and as little material as possible should be used, it is necessary to produce the mandrels close to the final contour during forging. To ensure this, it is necessary to estimate the probability of plastic deformation in the die. For this reason, the equivalent stresses are analysed in this context. Due to the arrangement of a reinforcement ring around the shoulder insert, the material can be prevented from plastic deformation. Nevertheless, the aim of this process should be to minimise the load on the die sections.

<figure><img src="image_790.png" /><figcaption>Fig. 3. Result of the numerical process design (a) and the die load at 5 mm Inconel thickness, 10 mm/s press speed and 1,000 °C furnace temperature (b).</figcaption></figure>

Subsequently, experimental forging tests were performed. During the tests a maximum press force of 1,960-2,000 kN was reached. A nearly similar maximum force of 2,050 kN was calculated by numerical simulation. Overall, the experimental and numerical force-displacement curves show good agreement (Fig. 4a). The characteristic ranges and the maximum force are correctly depicted in the simulation. Deviations in the force curve at high stroke values can be attributed to elastic deflection of the press system in reality. Further, a geometric verification was carried out. For this purpose, the experimentally produced die was optically measured and a surface file was exported. A comparison between experimental measured and numerical calculated die geometry is shown in