Tensile and Fatigue Performance of Cold-Work Tool Steels for Adjustable Forming Tools

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The results of this study support the implementation of adaptive tools in high-volume metal forming processes. The findings are relevant for designers of adjustable forming tools that rely on controlled elastic deformation, such as adjustable-diameter ironing punches, because the generated tensile and fatigue data enable informed material and coating selection, improved tool life prediction, and optimization of tool geometry.

Abstract

Forming tools adjustable by tensile elastic deformations offer opportunities for improved process control and reduced wear in high-volume metal forming processes such as ironing. However, the lack of tensile and fatigue data for hardened cold-work tool steels limits their broader adoption. This study investigates the mechanical performance of three tool steels—Vanadis^®4 Extra SuperClean, Vancron^® SuperClean, and Caldie^®—through uniaxial tensile and fatigue testing, supplemented by destructive static and fatigue/wear tests on specimens representative of an adjustable ironing punch. Non-coated specimens exhibited ultimate tensile strengths above 2700 MPa with approximately 2% plastic strain, while coated specimens fractured in a brittle manner between 1600–1900 MPa. Fatigue life at stress ranges between 1450–1750 MPa varied from several thousand to over four million cycles, with crack initiation linked to non-metallic inclusions and precipitates 10–30 μm in size. Finite element simulations accurately linked failure observed in uniaxial tests to the component-level tests, confirming that first principal stress is a reliable predictor for punch failure. All punch specimens withstood 10⁶ cycles at diameter changes up to 140 μm (4‰), with coated punches exhibiting minimal wear and non-coated ones showing localized surface damage. The findings support material and coating selection for adjustable forming tools and highlight opportunities for further optimization.

1. Introduction

Tool steels are heat-treatable, highly alloyed steels widely used in the production of critical components for various manufacturing tools. These include active tool parts, such as moulds and inserts for injection moulding, punches, dies, and blank holders for metal forming, and passive components like stress rings and load plates. In metal forming, active tool parts—those in direct contact with the workpiece—typically experience compressive stresses. However, several operations such as cold forging, deep drawing, and ironing also introduce substantial tensile stresses, elevating the risk of cracking or chipping. Prestressing using stress rings or stripwound containers is widely deployed to mitigate tensile stress peaks in conventional tooling [1].

Although static tensile properties of tool steels have been studied for over half a century, the literature remains fragmented, with few studies taking a holistic approach applicable to practical tool design. Early work by Grønbæk [2] established empirical relationships between compressive and tensile strength for conventional grades, showing a pronounced reduction in tensile strength with increasing hardness. Marsoner et al. [3] conducted direct uniaxial testing on powder-metallurgical (PM) steels, reporting ultimate tensile strengths up to 3.2 GPa and 0.2% yield strengths up to 2.9 GPa without necking, underscoring the need for appropriate specimen geometries. Furthermore, they stressed that bending-based methods, frequently used by manufacturers, tend to overestimate tensile strength for high-hardness materials. More recently, Morri et al. [4] showed that tensile strength, notch sensitivity, and fatigue performance in PM cold-work steels are sensitive to heat-treatment, tempering conditions, and resultant carbide morphology. Despite the industrial relevance of PM grades (e.g., Uddeholm’s Vanadis^® and Vancron^® [5], and Böhler’s MICROCLEAN [6] families), commercial datasheets rarely provide tensile or tensile fatigue data for PM tool steels, complicating reliable material selection for tensile-dominated applications.

Fatigue behavior has been explored more extensively, especially with the advent of ultrasonic methods enabling gigacycle (>10⁷ cycles) testing and evidence supporting the lack of a fatigue limit in metals [7]. Sohar et al. concluded that PM cold-work grades generally exhibit higher fatigue strength than conventionally produced steels [8], with studies demonstrating the beneficial role of compressive residual stresses and a load-amplitude-dependent transition from surface-initiated to bulk-initiated failure [9,10]. The crack initiation stage is often the life-controlling phase in very-high-cycle regimes [11], and inclusions strongly influence fatigue life scatter [12]. In PM cold-work steels with high vanadium content, cracks typically originate at non-metallic inclusions or carbide clusters of ≈10–30 μm in size, with higher cleanliness correlating with longer lives [13,14]. Comparable relationships between carbide morphology and fatigue crack initiation have also been reported for conventional tool steels [15]. An earlier low-cycle fatigue study on cold-forging tool materials by Brøndsted and Skov-Hansen [16] similarly emphasized the dominance of crack initiation processes and demonstrated the limited ductility of hardened tool steels under cyclic loading.

Surface integrity and near-surface residual stresses further influence fatigue behavior. Mechanical and thermal finishing methods (grinding, cutting, EDM) alter surface states and residual stresses, thereby affecting high-cycle performance [17,18]. Laser shock peening can introduce stabilising compressive residual stresses and improve fatigue resistance in tool steels [19]. Coatings add another layer of complexity: PVD multilayers and nitriding-assisted duplex systems can reduce friction and wear in forming tools, yet they may modify near-surface microstructure, influencing both crack initiation and growth under tensile-dominated cycling [20,21]. These findings collectively indicate that coating and surface effects must be understood alongside bulk microstructure when tools are subjected to cyclic tensile loads.

The emergence of tooling systems adjustable by controlled tensile elastic deformation introduces new material requirements. The adjustable-diameter ironing punch, proposed by the authors of [22,23,24], is a novel tool designed for rapid diameter adjustments during production in high-volume forming scenarios. This innovation offers two key benefits: (1) expanding the punch diameter to a controlled value before each stroke (“punch expansion” in Figure 1) allows control over the reduction ratio and thus part dimensions; and (2) relaxing tensile stresses at the bottom of the stroke (“punch contraction” in Figure 1) reduces wear by lowering [22,23] or even eliminating [24] the punch retraction force, thereby minimising contact between the punch and the workpiece. An elastic punch for backward can extrusion proposed by Lowrie and Ngaile is another example of adjustable tools [25,26]. Such tools demand cold-work tool steels capable of withstanding substantial cyclic tensile stresses without yielding or fracturing. While material selection for ironing dies and mandrels has been researched in [27], the literature does not provide systematic tensile and tensile fatigue data for punch material candidates—coated and non-coated—under loading modes representative of elastically adjustable tools. This gap restricts robust material selection, predictive modeling of tool life, and industrial adoption of the adjustable punch.

Figure 1. Working principle of the adjustable ironing punch.

This study addresses this gap by (1) generating uniaxial tensile and fatigue data for three commercial cold-work tool steels, including coated and non-coated conditions; (2) performing destructive static and cyclic testing on punch-like specimens that replicate the tensile-dominated stress state of an adjustable ironing punch; and (3) linking uniaxial results to component-level behavior via finite element analysis of the punch–mandrel system. Section 2 outlines the selected materials, specimen geometries, and experimental and numerical methods. Section 3 presents and discusses the results. Finally, the study is summarized and concluded in Section 4.

2. Materials, Specimens, and Methods

This section outlines the materials, specimen configurations, and procedures used throughout the study. Section 2.1 describes the selected tool steel grades, applied heat treatments, and specimen preparation for uniaxial tensile and fatigue tests. Section 2.2 introduces the punch specimens, detailing their design and manufacturing. Section 2.3 describes the experimental procedures, while Section 2.4 summarizes the used characterization techniques. Finally, Section 2.5 describes the finite element modeling approach used to correlate the component-level test data with uniaxial test results.

2.1. Material and Specimens for Uniaxial Tests

2.1.1. Tool Steel Grades, Heat Treatments, and Coating

Three cold-work tool steel grades, produced by Uddeholms AB (Hagfors, Sweden), were selected for uniaxial tests based on consultations with the manufacturer:

• Vanadis^®4 Extra SuperClean (V4E) [28]: a PM tool steel grade, selected due to its high abrasive wear resistance and purity.
• Vancron^® SuperClean (VSC) [29]: a PM grade offering high galling resistance, purity, high achievable hardness, and suitability for non-coated applications. Its high galling resistance in the non-coated state has been achieved by nitriding the metal already in the powder form, effectively integrating a surface coating into the bulk material [29].
• Caldie^® (CAL) [30]: a cold-work tool steel produced via the electroslag remelting (ESR) process, marketed for high wear resistance and ductility, and suitability for coated applications. It is the only grade among the three with tensile and fatigue data available in the manufacturer’s datasheet.

Manufacturer-provided chemical compositions of the selected grades are summarized in Table 1.

Table 1. Typical chemical composition of the cold-work tool steel grades used in this study.

Commercial Name	Abbr.	Alloying Elements (wt%)								Ref.
		C	Si	Mn	Cr	Mo	V	N	Fe
Vanadis®4 Extra SuperClean	V4E	1.4	0.4	0.4	4.7	3.5	3.7	-	bal.	[28]
Vancron® SuperClean	VSC	1.3	0.5	0.4	4.5	1.8	10	1.8	bal.	[29]
Caldie®	CAL	0.7	0.2	0.5	5.0	2.3	0.5	-	bal.	[30]

Uniaxial test specimen naming conventions are illustrated in Figure 2, and the six specimen types used are listed in Table 2. Four specimen types were made from V4E, with three heat treatments (HT2, HT3, HT4) to achieve target hardnesses of 62, 63, and 64 HRC, respectively (see Table 3). Some HT3-treated specimens were left non-coated, while others were coated with Duplex-VARIANTIC^® (DV) to study the effect of the coating on tensile properties. The commercially available DV coating combines plasma nitriding with a 4–6 μm thick multi-layer TiAlCN coating, commonly used for reducing the coefficient of friction and further improving galling resistance [31]. All V4E specimens with HT2 and HT4 were coated with DV. Caldie^® was used for one specimen type with HT1 (60 HRC, DV-coated), and Vancron^® was tested in one non-coated specimen type with HT3 (62 HRC, non-coated). All specimens were vacuum-heat-treated at 3–6 bar by an industrial partner. The number of tested specimens is shown in Table 4.

Figure 2. Naming convention for uniaxial test specimens used throughout this study.

Table 2. Description of uniaxial test specimen types with their nominal compressive ($R_{c0.2}$) and tensile ($R_{p0.2}$) yield strength and hardness values.

Specimen Type	Material	Heat Treatment	Coating	Hardness (HRC)	${\mathrm{R}}_{\mathrm{c}0.2}$ (MPa)	${\mathrm{R}}_{\mathrm{p}0.2}$ (MPa)	Ref.
CAL-HT1-C	CAL	HT1	DV	60	2350	1900	[30]
VSC-HT3-N	VSC	HT3	N/A	62	2750	Not specified	[29]
V4E-HT2-C	V4E	HT2	DV	62	2600 ± 150	Not specified	[28]
V4E-HT3-C	V4E	HT3	DV	63	2700 ± 150	Not specified	[28]
V4E-HT3-N	V4E	HT3	N/A	63	2700 ± 150	Not specified	[28]
V4E-HT4-C	V4E	HT4	DV	64	2850 ± 150	Not specified	[28]

Table 3. Specifications of heat treatments used throughout this study. All heat treatments were performed by vacuum heat treatment at 3–6 bar.

Specimen Type	Stress Relief		Austenitization		Tempering
	Temp. (°C)	Time (h)	Temp. (°C)	Time (min)	Temp. (°C)	Time (h)
HT1	650	2	1030	30	540	3 × 1
HT2	650	2	1030	30	525	3 × 1
HT3	650	2	1080	30	540	3 × 1
HT4	650	2	1150	15	525	3 × 1
HT5	650	2	1060	30	540	3 × 1

Table 4. Number of tensile and fatigue tests conducted with uniaxial test specimens, including the stress ranges ($\Delta \sigma$) used for the fatigue tests (30 Hz, R = 0).

Specimen Type	Tensile Tests	Fatigue Tests
		$\Delta \sigma$ (MPa)	n
CAL-HT1-C	3	1548	6
		1448	5
VSC-HT3-N	3	1692	1
		1648	2
		1598	1
		1548	3
		1498	1
		1448	1
V4E-HT3-N	3	-	-
V4E-HT2-C	3	1748	3
		1696	4
		1648	5
V4E-HT3-C	3	1728	2
		1695	3
		1660	1
V4E-HT4-C	3	1647	5
		1598	5

2.1.2. Specimen Geometry and Preparation

The geometries of uniaxial tensile and fatigue test specimens are shown in Figure 3. Both types featured a Ø6 mm gauge section and a Ø12 mm grip section. Tensile specimens were designed according to ISO 6892-1 [32], with a proportionality coefficient of k = 4.7 between the original gauge length and the square root of original cross-sectional area. ASTM E466-21 [33] was followed for the fatigue test specimen, with 0.5 mm deep grooves added to the grip sections to enhance gripping strength.

Figure 3. Manufacturing drawings of (a) ISO 6892-1 cylindrical uniaxial tensile test specimens and (b) ASTM E466-21 cylindrical uniaxial fatigue test specimens with added grooves for improved gripping in the test machine.

Specimens were manufactured by an industrial partner following the process route shown in Figure 4. The final polishing was performed on a RAP^® machine (Strecon A/S, Sønderborg, Denmark) with 9 μm diamond paste on an acrylic substrate to achieve a surface condition similar to comparable metal forming tools. A normal force of 10 N was applied during polishing. In the RAP^® machine, the specimen rotated at 600 min⁻¹, while the substrate oscillated axially at 75 Hz with a 1 mm range and moved along the specimen axis at 30 mm/min. Polishing was stopped once the axial line roughness, measured using a tactile profilometer by the manufacturer, reached the target value shown in Figure 3.

Figure 4. Process route for manufacturing of all uniaxial test specimens for this study. Coating was applied as the last step, where applicable.

Surface roughness of the gauge section was characterized for 24 out of 66 specimens in the as-received condition (before coating) using an S neox optical profilometer (Sensofar, Barcelona, Spain) in interferometric mode at $10\times$ magnification. The average axial Ra value was 0.043 μm with a standard deviation of 0.015 μm (n = 88), ranging from 0.025 μm to 0.085 μm. Eight of these specimens were also measured post-coating, and the DV coating was found to have no significant impact on the surface roughness. Due to the polishing mechanics of the RAP^® machine, the gauge section exhibited a repeatable hatched surface texture, as shown in Figure 5a. The texture direction map (Figure 5b) revealed two dominant directions between approximately $\pm {50}^{\circ }$ and $\pm {60}^{\circ }$ relative to the axial loading direction.

Figure 5. Typical surface condition of uniaxial tensile and fatigue test specimens: (a) gauge section surface texture after form removal and filtering; (b) texture direction.

2.2. Specimens for Adjustable Punch Tests

The geometry of the ironing punch test specimens is shown in Figure 6a. The thin-walled nose section (Detail A) closely resembles a previously presented tool in [24], but the thick-walled section was shortened to improve alignment and reduce cost. The nose section surface was finished by robot-assisted polishing as described in [24]. V4E with HT5 (see Table 3), resulting in a hardness of 62 HRC, was selected for these specimens. A total of 12 specimens were tested: 6 DV-coated and 6 non-coated. Half (3 coated and 3 non-coated) were used for destructive static tests, while the remaining were used for fatigue and wear testing.

Figure 6. Defined geometry of (a) the punch specimen and (b) the mandrel used to expand the punch specimen.

The mandrel used to induce deformation in the punch specimens is shown in Figure 6b. It was made from cemented carbide grade RG5, selected for its high stiffness [27]. RG5 has a mass density of 14,100 kg/m³, elastic modulus of 550 GPa, and hardness of 88 HRA (Silveralloy Co., Ltd., kasai, Japan, 8 October 2024). The punch–mandrel interface angle was 5°, consistent with previous studies [23,24,27].

2.3. Experimental Procedures

2.3.1. Uniaxial Tensile Tests

Uniaxial tensile tests were conducted using an Instron 8521 servo-hydraulic test machine [34] equipped with an MTS FlexTest 100 controller (Eden Prairie, MN, USA). Force was measured with an MTS 661.20E-03 load cell (±100 kN), and extension was recorded with an Instron 2620-602 strain gauge extensometer (Norwood, MA, USA) (25 mm initial gauge length, ±2.5 mm measuring range), as shown in Figure 7a. The tests were displacement-controlled at 0.2 mm/min and terminated upon specimen failure. To enhance grip strength and protect the machine’s hardened grips (60 HRC), 1.5 mm thick DC04 steel sheets were placed between the specimen and the grips, as illustrated in Figure 7b.

Figure 7. (a) Photo of a coated tensile test specimen in the test machine with the extensometer mounted. (b) Photo of a tensile test specimen fracture surface with the steel gripping sheets visible in the background.

Test data, including time, tensile force, and engineering strain, were recorded at 100 Hz. Post-processing was performed in Python 3.13, where data were averaged over 100 data points. The processing included calculating engineering stress, true stress, true strain, elastic modulus, and yield strength. The elastic modulus was determined as the slope of the true stress–true strain curve between 100 MPa and 1000 MPa. Yield strength was identified by the 0.2% offset method for specimens exhibiting sufficient plastic strain.

2.3.2. Uniaxial Fatigue Tests

Uniaxial fatigue tests were conducted using a servo-hydraulic test machine [35] equipped with an MTS 318.25 load frame, MTS FlexTest 100 controller, and an MTS 661.22D-01 load cell (±250 kN). As in the tensile tests, intermediate DC04 steel sheets were used between the specimen and grips to enhance clamping and prevent damage. The tests were force-controlled and performed at a frequency of 30 Hz with a load ratio of R = 0. The applied stress ranges for each specimen type are listed in Table 4. Peak and valley force values ($F_P$ and $F_V$) and machine displacement were recorded at each cycle. These data were used to calculate the stress range $\Delta \sigma$ according to Equation (1).

$$\Delta \sigma ={\displaystyle \frac{F_P-F_V}{A_i}}$$

(1)

2.3.3. Destructive Static Testing and Fatigue and Wear Testing of Punch Specimens

The most reliable fatigue life data for a component are obtained under actual service conditions. However, for the adjustable ironing punch, in situ testing would be costly and time-consuming, require significant development, and be potentially unsafe. Therefore, emulative tests, where the most relevant stress state for component failure occurs, were used. Two primary stress states occur during punch operation [24]:

• Post-expansion, pre-ironing: Dominated by hoop stresses in the thin-walled nose section, resulting in the highest first principal stresses.
• During ironing: External radial compressive forces counteract the tensile hoop stresses, producing the highest effective stress.

Due to limitations in applying radial compressive stresses reliably with the available equipment, fatigue testing was conducted under the first stress state (post-expansion, pre-ironing) where tensile stresses are highest. Consequently, the tests only represent the critical operational stress state, rather than all stress states. For production-ready tools, additional safety factors should be considered unless improved offline testing methods become available. All tests with the adjustable ironing punch specimens were performed on an MTS 312.21 servo-hydraulic test machine [36] with an MTS FlexTest 100 controller and an MTS 661.21A-03 load cell (±100 kN static capacity).

The test setup is shown in Figure 8. Two identical adapters interfaced with the machine grips and held the punch specimen and mandrel in place using a retainer ring. Radial deformation ($\Delta r$) of the punch was measured using three AXR/0.5/S LVDTs (Solartron Metrology, Bognor Regis, UK) with 1 mm range, spaced ${120}^{\circ }$ apart as shown in Figure 8. The change in punch specimen diameter ($\Delta d$) was calculated using Equation (2):

$$\Delta d=2·{\displaystyle \frac{\Delta r_1+\Delta r_2+\Delta r_3}{3}}$$

(2)

Figure 8. (a) Section view of the punch structural test setup with main dimensions. (b) Setup used for the destructive quasi-static and fatigue tests (lubrication equipment is not present in the photo) in the test machine.

Static tests were displacement-controlled at a mandrel velocity of 0.2 mm/min. The mandrel was inserted into the punch specimen (upward in Figure 8) until either failure or contact with the thick-walled section of the punch occurred. The conical contact surface was lubricated with Rhenus SU 500 A metal forming oil before the test, consistent with the actual application of the tool [24]. Mandrel position, axial force, and LVDT readings were logged at 100 Hz.

Fatigue tests were conducted at 5 Hz using sinusoidal, displacement-controlled mandrel motion (see Figure 9). Lubrication was maintained by continuously pumping Rhenus SU 500 A into the punch at 10 mL/h using a DME-8-10A dosing pump, representing realistic production conditions. A gap was allowed to form at the valley mandrel position ($x_V$), resulting in zero axial force ($F_V$) and diameter change ($\Delta d_V$), which facilitated lubricant replenishment at the contact interface. During the fatigue tests, peak and valley values of all relevant sensors were logged at each cycle.

Figure 9. Prescribed mandrel movement with the resulting axial force and specimen diameter change during punch specimen fatigue tests.

Tests with two different intended loading patterns were conducted on both coated and non-coated specimens. Two mandrels were used: one for all static tests and fatigue tests with non-coated specimens, and another for fatigue tests with coated specimens. The first type of test (Type 1) consisted of 100,000-cycle segments, the first of which was started at $\Delta d\approx$ 90 μm. The mandrel displacement amplitude was increased after every segment to achieve a diameter change 10 μm higher than at the beginning of the previous segment. After a total of 500,000 cycles, each subsequent segment was started at $\Delta d\approx$ 140 μm. The intended punch diameter change during the second type of test (Type 2) was constant at 140 μm. During the Type 2 test, a larger gap (≈20 mm) was introduced between the punch and the mandrel for 10 s after every 10,000 cycles to allow the lubricant to completely flush the contact interface. Both test types were stopped after wither specimen fracture or a total of ${10}^6$ loading cycles were reached. One coated and one non-coated specimen were tested using the Type 1 test, and two of each were tested using the Type 2 test.

2.4. Specimen Characterization

Fracture surfaces of uniaxial test specimens were examined using a Sigma 300 scanning electron microscope (SEM; Carl Zeiss IQS Deutschland GmbH, Oberkochen, Germany) equipped with an energy-dispersive X-ray spectroscopy (EDS) detector. Imaging was performed at 10 kV accelerating voltage with a 30 μm aperture. Elemental analysis was conducted at 20 kV with a 120 μm aperture size.

Wear development at the conical contact interface during punch fatigue and wear tests was monitored using an LEXT OLS4100 laser microscope (Olympus Corporation, Hamburg, Germany). Scans of approximately $2\times 2$ mm² areas were acquired at selected areas of interest on the mandrel. Post-processing was performed using MountainsMAP^® Premium (version 9.3.10249), involving removal of the conical form using a best-fit third-degree polynomial surface, outlier elimination, and application of a 0.8 μm Gaussian S-filter. A surface scan of the conical contact surface on the mandrel in the as-received condition is shown in Figure 10. The ISO 25178 [37] surface roughness parameters were Sa = 0.025 μm and Sq = 0.032 μm. Due to geometric constraints, the punch-side contact surface could not be scanned with the laser microscope. Instead, qualitative analysis was performed using photographic documentation.

Figure 10. LEXT laser microscope scans of the conical contact surface of the mandrel acquired at $20\times$ magnification showing (a) reflection intensity map and (b) topographical information. See “punch–mandrel contact surface” in Figure 8 for location of the scan.

2.5. Finite Element Modeling

The stress state in a uniaxial test specimen differs significantly from that in a punch specimen during the component-level test. To bridge this gap and enhance understanding of hardened cold-work tool steels under tensile loading, finite element simulation of the static punch test, described in Section 2.3.3, was conducted.

The commercial finite element software LS-DYNA was used to simulate the test. Given the axisymmetric nature of the punch and the mandrel, volume-weighted 2D axisymmetric solid elements were used. The simulation used a dynamic explicit time integration scheme, with both automatic mass scaling and velocity scaling applied to optimize simulation time. The minimum timestep size for automatic mass scaling was set to $2.25\times {10}^{-8}$ s. Mesh geometries and details are illustrated in Figure 11. Boundary conditions were defined as follows: the punch specimen was axially (z) constrained at the flange, while the mandrel was prescribed to move in the +z direction at 1000 mm/s via nodes under its flange. Contact between the punch and the mandrel was modeled using the CONTACT_2D_AUTOMATIC_ SURFACE_TO_SURFACE keyword with a static friction coefficient of 0.011 to maintain consistency with [24]. Both tool materials were modeled as linearly elastic, using the parameters shown in Figure 11. Simulation outputs included forces, stresses, nodal displacements, and contact forces, which were extracted for further analysis. A similar model was validated in [24] against process forces and final part geometry from experiments, so separate validation was not pursued in this study.

Figure 11. LS-DYNA simulation setup for numerical analysis of the stress state in the punch specimens using 2D axisymmetric solid elements.

3. Results and Discussion

3.1. Uniaxial Tensile Tests

Stress–strain curves from all 18 tensile tests are presented in Figure 12, along with average values for elastic modulus (E), 0.2% yield strength (${\sigma }_{y0.2}$), ultimate tensile strength (${\sigma }_u$), and fracture strain (${\epsilon }_u$). The results showed high repeatability, with most curves overlapping.

Figure 12. Stress–strain curves of uniaxial tensile tests (0.2 mm/min) of (a) CAL-HT1-C, (b) V4E-HT3-N, (c) VSC-HT3-N, (d) V4E-HT4-C, (e) V4E-HT3-C, and (f) V4E-HT2-C specimen types with three repetition averages of elastic modulus (E), 0.2% yield strength (${\sigma }_{y0.2}$, where applicable), ultimate tensile strength (highest engineering stress, ${\sigma }_u$), and fracture strain (${\epsilon }_u$).

Non-coated specimens, shown in Figure 12b,c, fractured after about 2% plastic strain before reaching plastic instability at an ultimate tensile strength above 2700 MPa with ${\sigma }_{y0.2}$ above 2250 MPa. In contrast, coated specimens (Figure 12a,d–f) fractured in a brittle manner, without significant plastic deformation between σ_u = 1600 MPa and 1850 MPa, depending on the specimen type.

A direct comparison between V4E-HT3-N (non-coated, Figure 12b) and V4E-HT3-C (coated, Figure 12e) confirms that the DV coating significantly reduces ductility. Despite the low toughness, coated specimens showed consistent fracture strains between 0.87% and 0.98%, with minimal scatter. The datasheet of Caldie^® specifies ${\sigma }_{y0.2}=1900$ MPa and ${\sigma }_u=2500$ MPa with fracture strain reaching ${\epsilon }_u=3\%$ (hardness unspecified) [30]. The CAL-HT1-C specimens underperformed significantly relative to these values, with the fracture strength (1662 MPa) being about 13% lower than the specified yield strength.

The DV coating combines plasma nitriding—which induces compressive surface stresses—and a multi-layer TiAlCN PVD coating, yielding a surface hardness of 3500 ± 500 HV_0.05 [31]. While plasma nitriding is known to enhance fatigue life by inhibiting crack initiation [38], it can also lead to surface embrittlement that can be detrimental to ductility. Alternatively, the coating application temperature of up to 500 °C [31] may have altered the microstructure, reducing ductility. Tekin et al. [39] observed ductility reduction in additively manufactured maraging steels coated with duplex coatings similar to DV. They explained the phenomenon with the brittleness of the coating and the diffusion layer that give rise to initial cracks, indicating that both plasma nitriding and PVD coating may contribute to the ductility reduction. The transferability of their findings to the present work remains unclear due to the material grade differences, so future tests isolating the effects of plasma nitriding and PVD coating could help explain the phenomenon in more detail.

3.2. Uniaxial Fatigue Tests

All uniaxial fatigue test results are shown in Figure 13. The measured stress ranges were approximately normally distributed around the target values, so variations can be considered random. Differences in error bar sizes may be due to servo valve wear affecting machine control.

Figure 13. Uniaxial fatigue test (30 Hz, R = 0) results showing number of cycles until failure at different stress ranges for (a) V4E-HT4-C, (b) V4E-HT3-C, (c) V4E-HT2-C, (d) CAL-HT1-C, and (e) VSC-HT3-N. Error bars show $\pm 2$ standard deviations of the stress range distribution. The three specimens highlighted using colorful circles in (c) are presented in Figure 14.

Among the coated V4E specimens, fatigue life varied significantly at constant stress ranges. Three V4E-HT4-C specimens (Figure 13a) failed after less than 25,000 cycles at Δσ = 1648 MPa, while one specimen fractured after about 617,000 cycles, and another one after 2.79 × 10⁶ cycles. Significant variation was observed for other specimen types as well with fatigue life at a constant stress range varying up to 900,000 cycles for V4E-HT3-C (Figure 13b), V4E-HT2-C (Figure 13c), and CAL-HT1-C (Figure 13d). The tested VSC-HT3-N specimens showed the least variation as seen in Figure 13e, but fewer repetitions were performed with this specimen type.

These results suggest that the statistical significance of the fatigue results is insufficient to demonstrate minor effects, such as 1-2 HRC hardness differences due to heat treatment. Large scatter in the fatigue life of high-strength tool steels is a known phenomenon, especially for high-cycle fatigue where crack initiation dominates variability, as observed by Kazymyrovych et al. [11].

Comparing coated V4E specimens with different heat treatments reveals that V4E-HT4-C (highest hardness, Figure 13a) exhibited fatigue strength about 50–100 MPa lower than HT2 and HT3 variants. This aligns with the inverse relationship between hardness and toughness, where reduced toughness may lower resistance to crack growth. However, Brønsted and Skov-Hansen [16] stated that fatigue life in PM tool steels is governed by crack initiation, instead, contradicting the traditional explanation. The difference between the fatigue life of V4E-HT3-C (optimal toughness and hardness, Figure 13b) and V4E-HT2-C (highest toughness, Figure 13c) was smaller than the difference with V4E-HT4-C (Figure 13a). Interestingly, V4E-HT2-C specimens tested at Δσ = 1696 MPa outperformed those tested at 1648 MPa, contradicting typical S-N curve behavior. However, due to the limited number of repetitions, this observation may be attributed to experimental variation. CAL-HT1-C specimens (Figure 13d) showed fatigue strength 100–200 MPa (around 10%) lower than other coated specimens, consistent with their lower static fracture stress (Figure 12), and lower hardness.

Figure 14 shows fracture surfaces of three V4E-HT2-C fatigue test specimens. Crack initiation sites (white arrows) were categorized as surface-initiated and bulk-initiated, and this classification is reflected in Figure 12 and Figure 13. Since this information was gathered after initial metallographic investigations, the crack initiation location of some specimens could not be retrieved. Notably, surface-initiated cracks occurred only at the highest stress ranges and always before 30,000 loading cycles, while bulk-initiated failures were dominant for specimens tested at lower stress ranges, reaching a greater number of cycles. In contrast, all V4E tensile test specimens (coated and non-coated) with a known fracture initiation site failed via surface-initiated cracks (Figure 12). According to Sohar et al. [9,10], bulk-initiated cracks are typical for specimens with high compressive residual surface stresses. Although surface stress characterization was outside the scope of this study, plasma nitriding and machining are known to introduce compressive residual surface stresses. Sohar et al. [9] also noted that at high stress amplitudes, the majority of residual stresses are released during the first loading cycle, increasing the likelihood of surface crack initiation. This explains surface-initiated failures observed in both tensile and high-stress fatigue tests.

Figure 14. Fracture surfaces of V4E-HT2-C fatigue test specimens failed after (a) 372,488 cycles, (b) 712,511 cycles, and (c) 28,227 cycles. The cracks for specimens in (a,b) initiated in the center of the specimen, while the crack in (c) was initiated at the surface of the specimen. The specimens are also encircled in Figure 12c with matching picture border colors.

Bulk-initiated crack sites were further studied using SEM-EDS. Secondary electron images of representative crack origins are shown in Figure 15. Particles approximately 10–30 μm in size were found at all initiation sites, consistent with the observations by Yao et al. [13]. In V4E-HT3-C (Figure 15a) and CAL-HT1-C (Figure 15c), the initiation sites showed evidence of fractured and pulled-out particles. EDS analysis of similar particles in inspected V4E-HT3-C (Figure 15a) specimens revealed the presence of oxygen, silicon, calcium, and aluminium, aligning with the typical composition of non-metallic inclusions found in tool steels [9,10]. A spectrum of the particle in Figure 15c could not be obtained due to its depth within the matrix.

Figure 15. Secondary electron SEM micrographs of crack initiation sites in (a) a V4E-HT3-C specimen, (b) a VSC-HT3-N specimen, and (c) CAL-HT1-C specimen. The CAL-HT1-C crack initiation site is shown at lower magnification in (d).

In VSC-HT3-N (Figure 15b), a larger particle was observed with a distinct morphology compared to the other two specimen types. Elemental analysis of VSC-HT3-N (Figure 15b) revealed that this particle consisted primarily of vanadium, nitrogen, and carbon, indicating that this could be a vanadium carbonitride (VCN) precipitate. These precipitates are a feature of Vancron^® SuperClean that enhance galling resistance, but the particle observed in Figure 15b is about one order of magnitude larger than expected [40]. The high hardness of VCN particles is beneficial for abrasive wear resistance, but it is detrimental for ductility and, when increasing to the size shown in Figure 15b, can act as a crack initiation site.

The fracture surfaces of V4E and VSC specimens exhibited traces of numerous small (around 1–2 μm) round particles. Similar surface morphology around the initiation site was also observed in [8], and it is consistent with the typical size and distribution of primary carbides. The smoother, circular region (around 200 μm radius) around the inclusion in CAL-HT1-C (Figure 15d) lacks signs of these primary carbides, but matches the typical fish-eye type crack origin described in [10,13].

In addition to non-metallic inclusions and precipitates, carbides can also act as crack initiators, especially if they have grown in size [14]. Processing steps with significant heat input, i.e., application of coating or surface treatment, could contribute to carbide growth, thereby increasing the number of potential weak points in the material. Although PM tool steels generally have a uniform chemical composition, such processes could induce segregation, which could also be detrimental to the fatigue properties of the material.

The crack initiation sites of V4E and CAL specimens differ fundamentally from those of VSC, since non-metallic inclusions were observed in the former and precipitates in the latter. As VSC-HT3-N was the only non-coated specimen type tested in uniaxial fatigue, the effect of the DV coating on fatigue strength could not be determined. The relatively low spread and low sensitivity to the applied stress range (Figure 13e) could thus either be caused by the lack of coating or the difference in particles found at the crack initiation site.

It is noteworthy that no delamination of the DV coating occurred in the gauge sections of any tensile or fatigue specimens, aligning with the observations of Bobzin et al. [21] and indicating that the tested material–coating combinations are suitable for further application-level testing.

3.3. Finite Element Modeling of Static Punch Specimen Tests

Simulated distributions of effective stress (${\sigma }_{eff}$) and first principal stress (${\sigma }_1$) in the punch and mandrel are shown in Figure 16a for diameter changes up to $\Delta d$ = 300 μm. The first principal stress increased linearly at 7.15 MPa/μm, closely matching prior analysis in [24] and thereby confirming that the test setup accurately replicates the relevant stress state for the punch.

Figure 16. (a) Simulated peak values of first principal stress ${\sigma }_1$ and effective stress ${\sigma }_{eff}$ (von Mises) at different punch specimen diameter changes. (b) Simulated first principal stress and (c) effective stress distribution in the punch nose at $\Delta d$ = 200 μm.

The distributions of first principal stress and effective stress in the mandrel and punch are shown in Figure 16b,c, respectively, for $\Delta d$ = 200 μm. The first principal stress in the punch peaks at the area above the punch nose and near the tip of the punch close to the interface and is governed by hoop stresses. Effective stress in the punch specimen peaks near the transition between the conical and cylindrical sections and is governed by compressive radial (r) and compressive axial (z) components, while the hoop stress is relatively low.

Contact pressure at the punch–mandrel interface, a key factor for wear [41], is shown in Figure 17. The peak contact pressure consistently occurs at around 6 mm from the punch tip, coinciding with the location of the maximum effective stress (Figure 16c) and increasing linearly from 550 MPa at $\Delta d$ = 50 μm to 2 GPa at $\Delta d$ = 300 μm at a linear rate of about 5.7 MPa/μm ($R^2=0.994$). The peak location shifted slightly (about 0.25 mm) along the z-axis with increasing deformation.

Figure 17. Simulated contact pressure distribution at the conical contact interface at specimen diameter changes up to $\Delta d$ = 300 μm.

3.4. Destructive Static Punch Specimen Test Results

Figure 18 presents the measured and normalized specimen diameter change and mandrel force during destructive static tests of the punch specimens. Non-coated specimens did not fracture within the test limits but exhibited plastic deformation before termination of the test. Coated specimens fractured at an average $\Delta d$ = 260 μm (7.1‰) diameter change without plastic deformation, similar to the V4E-HT3-C tensile specimens (Figure 12e). Non-coated specimens began yielding at about $\Delta d$ = 300 μm, setting a practical upper limit for low-cycle applications of the tool. Extrapolating from tensile test data (Figure 12b), non-coated punch specimens are expected to fail between $\Delta d$ = 850–1000 μm.

Figure 18. Measured diameter change $\Delta d$ as a function of axial mandrel force during destructive static tests of non-coated and DV-coated punch specimens made from V4E with HT5.

For coated specimens, the average diameter change at fracture ($\Delta d$ = 260 μm) corresponds to ${\sigma }_{eff}\approx$ 2520 MPa and ${\sigma }_1\approx$ 1860 MPa (from Figure 16a). This aligns closely with the uniaxial fracture stress of V4E-HT3-C (${\sigma }_u$ = 1833 MPa, see Figure 12e), suggesting that first principal stress is a more accurate predictor of punch failure than effective stress in the complex stress state occurring in the component-level tests. The 0.2% yield strain from uniaxial tests (Figure 12b) corresponds to $\Delta d$ = 315 μm, which matches the onset of nonlinearity in Figure 18. Conversely, the observed plastic deformation at $\Delta d$ = 300 μm corresponds to ${\sigma }_1\approx$ 2145 MPa according to Figure 16a.

3.5. Fatigue and Wear Test Results

None of the six punch specimens failed during fatigue and wear testing up to ${10}^6$ cycles, following the two loading schemes described in Section 2.3.3. Figure 19 shows the specimen diameter change ($\Delta d$), calculated using Equation (2), for the three coated specimens. The data of Specimen 1 from a Type 1 test (Figure 19a) shows a decrease in $\Delta d$ by ≈40 μm every ${10}^5$ cycles, with the rate of change diminishing over time. After restarting the test a few hours later, the pattern repeated. While the Type 2 test with Specimen 2 (Figure 19b) started at $\Delta d$ = 140 μm (${\sigma }_1\approx$ 1000 MPa, 3.99‰), the diameter change settled at around $\Delta d\approx$ 100 μm (${\sigma }_1\approx$ 715 MPa, 2.85‰) after about 150,000 cycles. That decrease was compensated for in the second Type 2 test with Specimen 3 (Figure 19c) by adjusting the mandrel displacement amplitude after 50,000 cycles, thereby resetting $\Delta d$ to 140 μm (3.99‰) and achieving a constant diameter change amplitude for the remainder of the test.

Figure 19. Measured peak diameter change of the punch specimens during fatigue and wear testing of coated (a) Type 1 and (b) Type 2 specimens at 5 Hz. Figure (c) shows the test data from the second Type 2 test, where the mandrel displacement amplitude was readjusted after 50,000 cycles.

The mandrel position, measured with the test machine’s built-in displacement sensor, was the control variable and was kept constant throughout the test. The observed decrease in $\Delta d$ over time—repeating only after long pauses—suggests influence from test machine temperature. Although machine temperature was not continuously monitored, spot measurements revealed that the temperature of the bottom grip reached 40–50 °C. This behavior could be explained by thermal expansion, which may have caused the actual mandrel displacement to decrease relative to the built-in LVDT reading, thereby reducing the effective interference and $\Delta d$.

Photos of the conical contact surface of the punch specimens tested using the Type 1 test are shown in Figure 20. For the non-coated punch specimen, visible changes developed during the ${10}^6$ loading cycles. The affected area, highlighted by the red ellipse in Figure 20a, features scratch-like marks aligned with the mandrel movement direction. These marks were confined to the transition zone between the conical and cylindrical surfaces—consistent with the location of the contact pressure peak shown in Figure 17. No visible marks were observed on the coated punch specimen (Figure 20b).

Figure 20. Photos of the conical contact surface of (a) non-coated and (b) coated punch specimens after ${10}^6$ cycles of the Type 1 fatigue and wear test. A wear scar is visible on the non-coated specimen (in the red ellipse).

To further investigate the altered areas, surface scans of the corresponding mandrel areas were acquired and are shown in Figure 21. Although the reflection intensity maps indicated a more extensive discoloration on the mandrel used with coated specimens, visual inspection revealed no significant difference in the total affected area. Comparing Figure 21a–f reveals no significant and consistent increase in affected area size after the first ${10}^5$ loading cycles.

Figure 21. LEXT laser microscope reflection intensity maps of selected visibly changed areas on the conical contact surface of the mandrel after (a,d) ${10}^5$ cycles, (b,e) $5\times {10}^5$ cycles, and (c,f) ${10}^6$ cycles used for fatigue/wear testing of (a–c) non-coated and (d–f) DV-coated V4E punch specimens. The scans were stitched from $3\times 3$ acquisitions taken at $20\times$ magnification.

Topographical maps of the same areas are presented in Figure 22. For the mandrel used with the coated punches (Figure 22d–f), discoloration did not significantly alter the surface topography when compared with the reflection intensity maps (Figure 21d–f). In contrast, scratches aligned with mandrel movement were clearly visible in Figure 22a–c, even in areas without visible discoloration in Figure 21a–c. This suggests that the discoloration could be due to lubricant polymerization under high contact pressure, rather than material transfer or removal. However, the scratches on the mandrel (Figure 22a–c) indicate material removal due to abrasive wear. Although the wear mechanism on the punch side could not be confirmed due to surface scanning limitations, the high hardness of the punch material suggests that the wear on the non-coated punch (Figure 20a) is mainly abrasive in nature.

Figure 22. LEXT laser microscope topographical maps of the same visibly changed areas shown in Figure 21 on the conical contact surface of the mandrel after (a,d) ${10}^5$ cycles, (b,e) $5\times {10}^5$ cycles, and (c,f) ${10}^6$ cycles used for fatigue/wear testing of (a–c) non-coated and (d–f) DV-coated V4E punch specimens. The scans were stitched from $3\times 3$ acquisitions taken with $20\times$ magnification.

The mandrel used for testing the non-coated punches had also been used in all six static tests, meaning minor pre-existing wear may have contributed to earlier wear initiation in the first (Type 1) fatigue test. However, both Type 2 tests were conducted with mandrels already subjected to at least ${10}^6$ cycles, and no significant difference in wear was observed. Notably, none of the coated punch specimens exhibited visible surface changes after 10⁶ cycles, while all non-coated specimens showed clear signs of wear. The substantially higher surface hardness of DV-coated specimens (3500 ± 500 HV_0.05 [31]) compared to non-coated specimens is likely the primary factor contributing to the observed reduction in wear.

The adjustable ironing punch is one of the few documented tools that rely on tensile elastic deformation for active adjustment in metal forming. Previous performance studies of that tool used two axisymmetric workpiece geometries with internal diameters between 33–35 mm and initial blank thicknesses of 1 mm [22,23] and 0.28 mm [24]. In those studies, the punch diameter change was limited to $\Delta d$ = 78 μm (2.31‰) [23], and $\Delta d$ = 120 μm (3.42‰) [24]. The present study shows that the tool could safely operate at diameter changes up $\Delta d$ = 260 μm in low-cycle conditions without external loading. In ironing applications, where radial compressive loads are present, lower limits should be applied. Although these loads reduce tensile hoop stresses, they increase the compressive radial stresses that govern the effective stress in the punch. Since the component-level tests here do not replicate all operational stress states, application-specific upper limits of diameter change should be validated separately.

Given that cemented carbide dies and mandrels reduce the required punch expansion [27], and that contact between the punch and workpiece was lost after retraction at $\Delta d>$ 90 μm (2.56‰) in [24], the practical adjustment range for process control lies between $\Delta d$ = 90–210 μm (2.56–5.98‰). This indicates that the novel tool could be capable of compensating for strip thickness variations as large as 60 μm for workpieces around 35 mm in diameter. This capability is sufficient to compensate for standard strip thickness tolerances of up to 1 mm in nominal thickness for the “normal” tolerance class, up to 1.5 mm for “fine”, and up to 3 mm for continuously rolled “precision” tolerance strips under 125 mm in width [42].

Wear at the punch–mandrel interface in non-coated specimens could be reduced by optimizing the contact geometry to reduce peak contact pressure while maintaining full contact across the adjustment range. This was also identified in [24] as a potential improvement to reduce mandrel push-out during ironing.

3.6. Material Selection and Scalability

Among the studied grades, Vanadis^®4 Extra SuperClean was selected for the adjustable forming tool due to its relatively high fatigue strength and wear resistance when coated with DV. Heat treatment variations resulted in minor differences in uniaxial tests, indicating that heat treatments achieving a balance between high hardness and toughness are beneficial. For cases, where high diameter change of the tool is not needed, coated Caldie^® could be a viable lower cost alternative. Vancron^® SuperClean could also be used for tools adjustable by tensile elastic deformations, but uniaxial fatigue test results (Figure 13) suggest that a lower diameter change should be used to avoid premature fatigue failure.

The achievable diameter change in adjustable forming tools is limited by their fatigue strength. Robust and safe operation of adjustable tools can be ensured by applying a safety factor and operating them at a lower diameter change than that limited by their fatigue strength. Even though non-coated tools are more ductile (Figure 18), making them more resilient against accidental excessive loading, they exhibit wear more than coated tools (Figure 20). The development of scratches due to wear reduces the predictability of fatigue life, because scratches concentrate stresses and could trigger surface-initiated fatigue failure at lower peak stresses. This supports the use of a higher safety factor for adjustable tools if used without a hard coating.

The adjustable ironing punch is scalable to different part diameters and thicknesses. Geometrical and material strength constraints make downscaling the diameter more challenging than upscaling. When reducing the adjustable punch diameter, the stress and strain induced in the punch nose with a unit increase in punch diameter also increase, effectively increasing the slope of the curves in Figure 16. Considering the strain in the punch nose equal to the relative diameter increase for elastic deformations, the minimum punch diameter $d_{0,min}^{punch}$ can be estimated using Equation (3) based on required punch diameter change $\Delta d_{req}$ for the process, the maximum permissible first principal stress ${\sigma }_{1,max}$ and the elastic modulus of the punch material. As the ironed part diameter is essentially equal to the punch diameter, this equation can be used to estimate whether developing an adjustable-diameter punch is technically feasible for the specific part. Assuming E = 210 GPa and ${\sigma }_{1,max}$ = 1500 MPa, which are realistic values for hardened PM tool steels, the punch diameter should exceed 14 mm if at least 100 μm of punch diameter change is needed.

$$d_{0,min}^{punch}\approx {\displaystyle \frac{\Delta d_{req}}{{\sigma }_{1,max}}}·E$$

(3)

Manufacturability of the punch and mandrel becomes more challenging as the diameter decreases due to lower bending stiffness of the parts. The low desired wall thickness of the punch nose and high fatigue strength are conflicting requirements, because the relative size of non-metallic inclusions, carbides, and precipitates, which act as stress concentrators, grows. Reducing wall thickness also increases the effect of any applied coatings and surface treatments, thereby making the tool vulnerable to fatigue failure, and reducing the predictability of failure.

4. Summary and Conclusions

Industrial implementation of forming tools adjustable by tensile elastic deformation benefits significantly from a deeper understanding of the fatigue performance of cold-work tool steels under tensile loading. This study addressed this knowledge gap by conducting

• Uniaxial tensile and fatigue tests on three Uddeholm cold-work tool steels at R = 0 and 30 Hz: two powder-metallurgical (PM) grades and one electroslag remelted (ESR) grade.
• Destructive static and fatigue/wear tests on adjustable-punch-like specimens actuated by a conical mandrel, simulating the stress state with the highest tensile loads in an adjustable ironing punch.

Experimental and numerical results support the following conclusions:

• Static tensile testing showed that non-coated V4E and VSC specimens sustained ≈ 2% plastic strain and fractured above 2700 MPa, whereas DV-coated specimens failed in a brittle manner between 1600–1900 MPa. This confirms that the coating, typically used under compressive loads, substantially reduces ductility while maintaining consistent fracture behavior.
• Uniaxial fatigue tests across all grades exhibited large scatter in fatigue life at stress ranges of 1450–1750 MPa, with failures occurring from a few thousand to several million cycles. High-stress tests primarily produced surface-initiated cracks, while specimens surviving beyond ≈30,000 cycles predominantly exhibited bulk-initiated failures.
• Bulk-initiated failures consistently traced back to inclusions or precipitates 10–30 μm in size, including large VCN precipitates in VSC and non-metallic inclusions in V4E and CAL. These findings highlight the critical influence of microstructural purity on fatigue performance.
• Destructive punch tests demonstrated that DV-coated punch specimens fractured at $\Delta d\approx$ 260 μm (7.4‰) without plasticity, while non-coated specimens yielded at $\Delta d\approx$ 300 μm (8.55‰) and did not fracture within the test limits. First principal stress correlated well with these observations, confirming its suitability as a failure predictor for adjustable tools operating in predominantly tensile stress states.
• All punch specimens—–both coated and non-coated—–survived ${10}^6$ cycles at diameter changes up to 140 μm (3.99‰), corresponding to first principal stresses of ≈1000 MPa. Wear was minimal on DV-coated specimens, while non-coated punches developed localized scratches at the peak-pressure zone.

This study provided valuable data on the tensile and fatigue performance of commercially available tool steels, where such information is often lacking. While all three studied grades could be used for forming tools adjustable through elastic deformation, higher safety factors should be selected for non-coated tools. Fatigue failures were initiated by inclusions and precipitates, highlighting the importance of microstructural homogeneity and purity of the material. The experiments also demonstrated the resilience of the DV coating under tensile loading and wear reduction at the punch–mandrel interface, supporting the use of hard coatings in adjustable forming tools.

Future research could investigate the mechanism behind toughness reduction due to the DV coating, potentially isolating the effects of plasma nitriding and PVD coating, and propose tool coatings tailored for use in forming applications under tensile loads. Furthermore, the punch specimens featured a conical contact interface geometry, which led to a non-uniform pressure distribution. Optimization of contact interface geometry could lead to a lower peak contact pressure, thus minimizing wear and prolonging tool life.