2.1 Material and methods

The punches were supplied in a TiN-coated state. The reported hardness was close to HRC 65. The applied heat treatment was quenched and tempered. A total of 7 HSS punches were used for this study. The coupons cut from the tip of these punches were used in the abrasion tests, and the ones cut next to them were used for characterization studies. The punches’ chemical composition and heat treatment history were received from the supplier. In addition, the chemical composition was verified with energy-dispersive X-ray spectroscopy (EDS) analyses conducted with low magnification. The heat treatment state was checked with hardness measurements. The specimen was verified prior to experimental tests. Table 1 shows the chemical composition of DIN 1.3343 steel. The steel contains strong carbide maker alloying elements such as W, Cr, Mo and V. This type of steel is suitable for nitriding as well as physical vapor deposition (PVD) coating.

Figure 1 shows the SEM image of used HSS steel. The microstructure contains small carbides distributed all over martensitic matrixes. The carbides were generally formed at grain boundaries.

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SEM image of supplied HSS steel

The punch used in this study is currently used to pierce these holes on the stainless steel sheet (AISI 430 with 0.6 mm thickness) where adjustment buttons such as temperature, time, and fan status, are located on the front panels of the ovens, as shown in Figure 2. In addition, since it is used in the production of final parts for assembly, no lubrication is performed. Because the final parts are taken directly from the production line to the assembly line, in this process, deburring and oil cleaning of the final parts create extra costs, and drilling problems seriously affect the scrap cost. However, the protective plastic strip could act as lubricant up to a level. The piercing speed was 8 holes/min. The punching process was carried out sequentially at room temperature in the form of piercing. Punches are changed in practice when the desired hole size cannot be produced or when the amount of burrs at the hole exit exceeds certain limits.

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The usage of used punch in present study

2.2 Heat treatment of punches

The traditional heat treatment applied to HSS steels typically begins with austenitizing at relatively high temperatures (1100-1300 ℃). The reason for performing austenitizing at elevated temperatures compared to conventional steels is to break down the carbides present in the structure and facilitate their dispersion throughout the matrix. Following this step, a quenching process is carried out to ensure the occurrence of the martensitic transformation, which is determined based on the steel’s composition. This process can be performed in air, water, oil, or special salt solutions, either in a single step or progressively. The aim is to prevent undesirable effects like distortion, brittleness, and crack formation that may result from rapid cooling. Subsequent to quenching, a tempering process is performed to reduce the brittleness of the martensitic structure. Low-temperature tempering (around 200 ℃) in this step results in a degree of toughness improvement while causing a slight reduction in hardness. If the alloy composition is sufficient, high-temperature tempering (above 500 ℃) can be applied. During this process, the formation of new carbides in the structure is encouraged, enhancing hardness through these new carbides while slightly reducing the hardness of the martensitic matrix. The high carbon content in the structure of HSS tools leads to the retention of a portion of austenite without transformation, referred to as retained austenite. While some of this austenite transforms during tempering, one of the procedures performed to induce the transformation of the remaining retained austenite and to generate secondary carbides within the structure is cryogenic treatment [19].

Cryogenic treatment is an additional heat treatment method applied to traditional heat treatments. The selection of two different tempering temperatures for the heat treatment procedures applied to punches is based on the assumption that carbides formed at higher temperatures could potentially contribute positively to wear resistance. In this study, the effects of cryogenic treatment and tempering applied to commercially available pre-fabricated punches were investigated. Table 2 summarizes the heat treatment procedure used in the paper. Based on the existing literature on the topic and the authors’ expertise, a deep cryogenic treatment at liquid nitrogen (-196 ℃) was selected for this study. The soaking durations were varied as 12, 24 and 36 h for the cryogenic treatment. Recognizing that high-temperature tempering and low-temperature tempering offer distinct advantages, both conditions were investigated, with tempering conducted at 200 and 500 ℃ for a duration of 2 h subsequent to the cryogenic treatment. Figure 3 shows the cryogenic treatment unit used in the study, and Figure 4 shows the time vs temperature graphs.

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Cryogenic treatment unit used in the tests

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Time-temperature cycle graphs

2.3 Wear tests

The wear tests aimed to simulate the accelerated version of the wear exhibited by the punches after a certain period of use in practical applications. This allowed for a controlled simulation of the punches’ wear under extended periods and varying conditions. For this reason, the wear tests were conducted in the same condition as they are used practically with TiN-coated surfaces. Samples taken from the cutting part of the punches were affixed using appropriate mounting fixtures and subjected to testing. The ball-on-disk wear test setup, known as the adhesive test setup, which is the most suitable for simulating the working conditions of the punches, was selected. The tests were conducted using a CSM tribometer unit, and the instantaneous friction force was measured and recorded during the tests. Measuring the friction force is particularly valuable for coated applications, as it helps determine at which stage the coating becomes damaged. The tests were conducted under a 10 N load, with a relatively small radius (1 mm) and a speed of 2.5 cm/s, over a distance of 100 m. The 10 N load was the maximum load allowed by the CSM tribometer. To increase the contact stress, a spherical ball with a diameter of 3 mm was used. The choice of using a very hard ball made of WC material was to ensure that wear only occurs in the tested material during the wear test. The selected wear radius and speeds were determined considering the working conditions. The procedure of specific wear rate calculation and details of equipment were presented in a previous study [20]. A sample wear section area measurement is given in Figure 5. The average wear area for each test sample was calculated as the mean value of many measured sections. The wear volume V_wear was obtained by revolving the average wear area A_av over the test radius per Eq. (1). The specific wear rate is a common approach that alters the wear test results by removing the influence of load and distance. The specific wear rate R_sw was calculated using Eq. (2).

(1)

(2)

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A sample wear area measurement for G0

where L is load and D is distance.

The practical punch used in this study has a significant life; in order to simulate its working conditions and obtain a reliable wear behavior in a reasonable time, high contact stresses are generated. For this purpose, a 10 N load was used on a 3 mm WC-ball. Ideally, if the ball and the specimens were rigid, the contact should be a point contact. The stresses exerted on and below the surface were calculated using Hertzian theory. The elastic modulus was taken at 200 GPa for tool steel, and the Poison ratio was taken at 0.3, while for WC-Co ball, the elastic modulus was taken at 690 GPa, and the Poison ratio was 0.22 per WC-Co ball specification. The contact stress was the same for all tested specimens. The test parameters are given in Table 3.

2.4 Microstructure and XRD analyses

Microstructural analysis was performed on the uncoated surfaces of the cut coupons from the punch. The HSS materials were gradually ground, polished, and etched to reveal their microstructure. Optical microscopy examination did not reveal significant differences among the samples, prompting the use of SEM. The SEM images were obtained at magnifications of 1000× and 5000× using a Hitachi Regilus 8230 model electron microscope to reveal detailed carbide structures. XRD measurements were conducted using equipment that covered an angle range of 10°-100° with Cu Kα radiation. Phase analysis was performed using XRD patterns; in addition, Rietveld analyses were carried out to determine the types and quantities of phases in the materials, using the material analysis using diffraction (MAUD) program. The analyses were continued until the sigma value, reflecting that the goodness of fit was the closest to 1.

3.1 Influence of cryogenic treatment on the wear resistance

Wear tests were conducted on the TiN-coated samples taken from the punch’s working region. This choice was made to examine the wear behavior under conditions resembling the practical application of the tool with the coating. In Figure 6, the friction coefficient vs test distance graph was given. The data were filtered to highlight the trend. The untreated punch (G0) showed the highest friction coefficient, while the samples soaked for 36 h (G3 and G6) showed the lowest friction coefficient. It is evident that as long as TiN coating maintains its integrity, it provides a lower friction coefficient of approximately 0.2-0.3 when in contact with the WC ball. However, as the coating wears off, a significant increase in friction coefficient is observed when the HSS material interacts with the WC ball. The observed friction increase in the G3 sample over 20 m of tests proves this theory, and as the test proceeds, the friction coefficients increase and come close to other samples.

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Friction coefficient vs distance

In Figure 7, the calculated specific wear rate ratios based on wear cross-sections are presented. According to the wear rates, it is observed that cryogenic treatment enhances wear resistance compared to the reference material in all conditions. The most optimal result is achieved with 36 h of deep cryogenic treatment followed by tempering at 200 ℃. Figure 7 shows the specific wear rate of specimens. The samples with the lowest friction coefficient (G3, G6) showed the highest wear resistance; since all the tested specimens have the TiN coating at the beginning of the test, it is believed that cryogenic treatment conducted at 36 h soaking time improved the wear resistance of tool steel which retarded the deterioration of TiN coating and presented an improved tool life. Figure 8 shows the selected wear scar sections measured with surface roughness tested. The illustrated sections are the ones that are the closest to the average wear area. As shown, comparing G0 to G3, the cryogenic treatment improves the wear resistance by retarding the failure of the coating resulting in a specific wear rate increase by a factor of 9 in the wear scar area.

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Specific wear rate of specimens

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Selected wear scar of specimens: (a) G0; (b) G1; (c) G2; (d) G3; (e) G4; (f) G5; (g) G6

3.2 Influence of cryogenic treatment on hardness

The effect of cryogenic treatment on hardness has been reported in the literature to be around HRC 1-2 on a macro scale. Macro hardness measurement methods provide more reliable results due to their broader scope. In this study, since the hardness of the HSS steels used was very close to the limits that could be measured macroscopically, hardness measurements were conducted on a microhardness scale using the Vickers method with a 0.3 N load. Figure 9 shows the graph of microhardness measurements. Cryogenic treatment and tempering resulted in a hardness increase of HV 20-30. Considering the method used, these increments are consistent but should not be considered significant.

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Microhardness measurements of samples (Vickers with 0.3 N)

3.3 Influence of cryogenic treatment on microstructure

The optical microscopy examination did not reveal significant differences between the samples subjected to cryogenic treatment and those not treated, although meaningful differences were only observable under SEM at 5000× magnification. Figure 10 displays the SEM images of specimens under 5000× magnification. When examining SEM images of specimens treated with cryogenic treatment compared to the reference (untreated) specimens, it can be noted that smaller secondary carbides were observed in the samples subjected to 36 h deep cryogenic treatment, which were not present in the reference specimens. Observing a similar effect for shorter soaking time is rather challenging. The applied cryogenic treatment is seen to result in the growth of carbide sizes when the tempering temperature after cryogenic treatment is increased from 200 to 500 ℃. Additionally, residual austenite in the matrix structure could not be distinguished through SEM examination. Consequently, XRD tests were performed to measure the amount of residual austenite in the structure and analyze the extent of the observed increase in secondary carbides.

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SEM images of HSS samples with different cryogenic treatment and tempering procedures: (a) G0; (b) G1; (c) G2; (d) G3; (e) G4; (f) G5; (g) G6

Figure 11 shows the energy dispersive X-ray spectroscopy (EDS) results of observable carbides (3-4 μm diameter). It is revealed that they are complex carbides M₆C that contain W, Mo, Fe, Cr, and V together. Figure 12 shows an example of MC carbides which contain mostly V and have smaller sizes (<1 μm).

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The EDS results of M₆C in G3

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The EDS results of MC in G3

XRD tests were conducted between 10° and 100° scanning angles, and the results obtained were initially analyzed using the Match software. In this software, the alloying elements present in the material’s structure were analyzed along with the information about the XRD source. These analyses were cross-referenced with previous studies found in the literature for validation. The XRD findings unveils that the reference microstructure consists of martensite, retained austenite, and MC and M₆C carbides. Figure 13 shows the peak location of each phase.

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XRD patterns of HSS samples

In the realm of HSS, the designations MC and M₆C are employed to denote specific carbide phases that form due to the interplay of alloying elements and heat treatment procedures during the steel’s fabrication. These carbide phases hold pivotal significance in dictating the steel’s hardness, wear resistance, and other mechanical attributes. MC carbides are carbides composed of a metal element (typically tungsten or molybdenum) and carbon. The letter M signifies the metal element (tungsten or molybdenum), while C represents carbon. These carbides are rigid and contribute to the overall hardness and wear resistance of HSS steels. They are generally more widely dispersed throughout the microstructure and contribute to the steel’s strengthening. On the other hand, M₆C carbides are intricate carbides generally comprising multiple elements, including chromium, tungsten, molybdenum, vanadium, and carbon. The M₆C designation indicates the involvement of six distinct elements (typically a combination of those mentioned) in shaping the carbide structure. M₆C carbides are harder and more stable at elevated temperatures compared to MC carbides. They confer heightened resistance to wear, high-temperature softening, and abrasion. The M₆C carbides generally form at higher temperatures and have larger dimensions compared to MC. They are also dispersed around microstructure, but they tend to form at grain boundaries. As the tempering temperature increases, there is a trend in the formation of more complex carbides. The specific types and distribution of carbides depend on the alloy composition of the HSS and the heat treatment processes that it undergoes. The equilibrium between these carbide phases, coupled with other factors such as grain size and matrix structure, dictates the comprehensive performance attributes of the steel, including hardness, toughness, wear resistance, and high-temperature stability. The tailoring of microstructure and heat treatment procedure was mostly made due to the working temperature. Tool steels such as 1.2379 HSS were not used at elevated temperatures, so the MC carbide, which has higher hardness at room temperature compared to complex carbides [21], is preferred.

Table 4 provides the phase percentages obtained through Rietveld analysis. The residual austenite phase, which is not easily discernible in SEM examination due to its distinct lattice structure compared to the martensitic phase, can be readily determined by using XRD analysis [22]. Fundamentally, the deep cryogenic treatment has reduced the residual austenite content in all durations and tempering temperatures, decreasing from an initial level of 7.81% to around 1%. The fact that the transformation from austenite to martensite involves a diffusionless transformation explains the proximity of the observed transformations across all samples. One of the mechanisms attributed to the effects of cryogenic treatment, namely the transformation of residual austenite to martensite, has been demonstrated through XRD tests and Rietveld analysis. The second widely accepted mechanism involves the contraction of the unit cell at cryogenic temperatures, leading to the diffusion of carbon within the lattice and promoting the formation of secondary carbides. Given the significant reduction in diffusion rates at cryogenic temperatures, observing the effects of this mechanism requires longer durations. The tempering process, especially at elevated temperatures, supports carbide formation when alloying elements conducive to carbide formation are present in the structure. This phenomenon is particularly observed in groups tempered at higher temperatures (G4, G5 and G6). Analyzing the impact of cryogenic treatment duration on MC and M₆C carbides in the structure would be more accurate when using samples tempered at lower temperatures. Upon examination of the trend, an increased cryogenic treatment duration has facilitated the transformation from simpler MC carbides to more complex M₆C carbides. The heightened presence of M₆C carbides in the structure corresponds to the observed increase in wear resistance.

In the context of HSS, the designations of MC and M₆C are employed to signify specific carbide phases that form due to the interplay of alloying elements and heat treatment during steel fabrication. These carbide phases play a significant role in influencing hardness, wear resistance, and other mechanical properties. MC carbides are composed of a metal element (commonly tungsten or molybdenum) and carbon, contributing to hardness and wear resistance. On the other hand, M₆C carbides consist of multiple elements, including chromium, tungsten, molybdenum, vanadium, and carbon, offering heightened resistance to wear and high-temperature effects. Both MC and M₆C carbides contribute to the overall properties of HSS steels, with specific types and distribution influenced by alloy composition and heat treatment.

Residual austenite, not easily identified in SEM, can be determined using XRD due to its unique lattice structure compared to martensite. Cryogenic treatment effectively reduces residual austenite in various durations and tempering temperatures. The transformation from austenite to martensite occurs via a diffusionless mechanism, resulting in consistent transformations across samples. One established mechanism of cryogenic treatment, the transformation of residual austenite to martensite, is confirmed through XRD and Rietveld analysis. In this study, a portion of the residual austenite remaining after cryogenic treatment transformed into martensite. Specifically, the retained austenite amounts, as determined through XRD tests and Rietveld analysis, underwent a transformation. The cryo-treated and tempered structure exhibited only a small quantity of retained austenite (ranging from 1.23% to 0.18%), which varied based on the cryo-treatment and tempering temperatures.

A second mechanism known as the formation of secondary carbides was observed in the present study. These duration varies with different alloys and under different initial heat treatment conditions. Notably, increased cryogenic treatment duration promotes transformation from simpler MC carbides to complex M₆C carbides, enhancing wear resistance. The impact of cryogenic treatment duration on MC and M₆C carbides is better analyzed through samples tempered at lower temperatures. The fundamental mechanism behind cryogenic treatment involves the contraction of the atomic lattice at low temperatures, causing some of the carbon within the structure to move out of the lattice. This carbon then combines with carbide-forming alloying elements to create secondary carbides. This phenomenon has been particularly noted with an increase in the M₆C carbide proportion within the structure. While the rise in tempering temperature to 500 ℃ after cryogenic treatment contributes to the increase in MC and M₆C carbides in the structure, it results in a slight decrease in the hardness of the martensitic matrix. Consequently, it offers only limited improvement in wear resistance despite the increase in carbide content compared to cryo-treated tempered specimens at 200 ℃.