Improving the Wear Resistance of Ledeburitic Tool Steels by a Combination of Semi-Solid and Cryogenic Processing

Abstract

Ledeburitic tool steel X210Cr12 was processed by passing through a semi-solid state with subsequent forming on a hydraulic press, thus achieving a fine-grained martensitic matrix, uniformly dispersed fine precipitates, and removal of sharp-edged primary chromium carbides. The hardness value was over 700 HV10. The decomposition of austenite and the promotion of further carbide precipitation were carried out by cryogenic treatment or a combination of cryogenic treatment and tempering. Transmission electron microscopy showed that tempering after cryogenic treatment also led to the precipitation of needle-like M₃C carbide, unlike the previous regimes. Furthermore, after the combined treatment, the microstructures showed a significant wear resistance, which was detected both by a waterjet abrasive blast test and a laboratory pin on disk test. Both tests showed a significant increase in wear resistance compared to the initial condition and special high wear resistance steels, such as Hardox 450 and Hardox 600.

1. Introduction

Abrasive wear is one of the basic damage modes of machine parts and components. Therefore, new materials, processing methods, or surface treatments are constantly being sought to obtain materials with higher wear resistance. Mostly, materials with a higher fraction of uniformly dispersed fine carbides in a martensitic matrix with a small fraction of metastable retained austenite are being sought [1,2]. Such materials are used in a wide range of applications, such as mining, metallurgy, road building, drilling, and other industries. One particular application may be for the replaceable wear liners (side liners, apron liners, side/frame liners) in the chambers of various types of horizontal shaft impactors [3]. These liners need to be replaced frequently due to abrasive wear caused by the impact of stones, concrete, construction material, or other material in the crusher chamber and during the crushing of the material by blow bars. They are usually made of AR (abrasion resistant) steel plates, which are high-strength martensitic steels [3,4]. To ensure the highest possible wear resistance, new materials [5] and new heat or thermomechanical processing methods, such as direct quenching and redistribution [6] as well as new manufacturing technologies [7,8,9], are being investigated.

In metal matrix-carbide composites, the hardness of the matrix and carbides, their distribution and size, and the interface strength between the matrix and carbides are important. For wear-resistant alloys based on Fe–Cr–C hardfacings, the size of the precipitating carbides and their resulting brittleness is an issue. Martensitic steels exhibit good impact abrasion resistance but have lower abrasion resistance due to the absence of hard phase precipitates. Complex Fe–Cr-based hard metals alloyed with W, Nb, B, Si, and other alloying elements are suitable for use in two- and three-body abrasive environments [5].

Another option is using conventional, high-alloyed ledeburitic tool steels processed by semi-solid state transition technology, which leads to a very hard martensitic matrix with finely precipitated chromium carbides with sizes of about 100 nm [7,8,9,10,11]. Semi-solid processing is a forming technology in which the material is partially melted. Due to the high temperature, the sharp-edged chromium carbides, which significantly reduce the material’s toughness, are dissolved. These M₇C₃ carbides are formed in the eutectic reaction during cooling and are typical for ledeburitic types of steels [12]. Due to the presence of the melt and with the appropriate size and shape of the solid particles, the material exhibits thixotropic behavior, which allows the production of shape-complex parts in one forming step to form even hard-to-form materials and to use lower forming forces than in classic forming technologies [13,14,15,16]. After semi-solid processing, the microstructure consists of polyhedral austenite grains surrounded by a ledeburitic network [17,18]. The grain refinement, the network morphology, and the promotion of fine carbide precipitation can be influenced by the cooling parameters and the forming conditions [19].

Since the microstructure after semi-solid processing contains a high fraction of austenite, it is possible to further increase the wear resistance by cryogenic treatment. Its purpose is mainly to stabilize the martensite and eliminate retained austenite by cooling the steel below the M_f temperature. This temperature is usually below −100 °C for tool steels. Cryogenic treatment (DCT) can be defined as an add-on process for hardening and tempering steels. In this process, the material is frozen to a temperature close to that of liquid nitrogen (−196 °C). Another reason for DCT, but even more important, is the stabilization of the martensite, which is accompanied by the precipitation of finer dispersed carbides during the tempering process. Tempering generally follows immediately after DCT. This process increases wear resistance in most tool steels [20,21]. DCT leads to the decomposition of retained austenite to martensite during cooling to temperatures down to −130 °C. During freezing to lower temperatures, i.e., down to −196 °C, the primary martensite decomposes during a time-dependent transformation leading to the nucleation of numerous coherent nanocarbides. The newly formed carbides are considered to be the main contributor to the significant increase in wear resistance after cryogenic treatment [20,21,22,23]. A higher quantity of precipitated nanocarbides can be expected with a longer holding time at cryogenic temperature. Increased tool life due to DCT has recently been described for various machining and forming tools [22,23].

Most of the research works dealing with wear resistance discuss the problems of various types of materials produced by standard metallurgical methods as well as by powder metallurgy and subsequently processed by conventional methods, such as heat treatment. The works on semi-solid processing deal mainly with the influence of processing parameters on the thixotropic behavior of the material, production of shape-complex parts, and obtaining a structure without internal defects. A combination of thermomechanical processing with a transition through a semi-solid state followed by cryogenic treatment to obtain materials with higher wear resistance has not been described so far. Therefore, this paper discusses the effect of combined semi-solid and cryogenic processing on the development of structures and their influence on wear resistance.

2. Materials and Methods

The experimental program was conducted on X210Cr12 tool steel with 1.8%C, 12%Cr, 0.3%Mn, and 0.35%Si (

Table 1. Chemical composition of X210Cr12 steel (wt.%).

C	Si	Mn	Cr	P	S
1.9–2.2	0.1–0.6	0.2–0.6	11–13	max. 0.03	max 0.01

). It is a ledeburitic steel with simple alloying without expensive alloying elements. The steel was supplied in the annealed condition, and the microstructure consisted of large sharp-edged chromium carbides and very fine cementite in a ferritic matrix (Figure 1a).

JMatPro (Sente Software Ltd., Guildford, UK), (Figure 1b) [24] was used to design suitable heating temperatures for the first experiments. The temperature of the start of liquid phase formation was determined to be 1247 °C. Complete melting of the material occurs at 1390 °C. The start of the formation of the austenitic phase was found at 722 °C. The dissolution temperature for all the carbides in the structure was also important and was found to be 1260 °C. For this steel, it was significant that the temperature interval between the solid and the liquid was 143 °C, giving a sufficiently wide temperature interval to obtain the necessary fraction of the liquid phase in the structure. Based on previous experience, the aim was to achieve approximately 30–40% liquid phase in the microstructure [11].

2.1. Semi-Solid Processing with Subsequent Thermomechanical Treatment

Semi-finished products with a diameter of 75 mm and a height of 105 mm were made from X210Cr12 and inserted into S235JR tubes with a melting temperature of 1400 °C and good formability. The ends of the tube were closed with caps of the same material and were welded around the circumference (Figure 2a). This container was created to safely handle the molten semi-finished product from the ledeburitic tool steel.

Since the tool steel semi-finished product is placed in a low-carbon steel container, other factors (heat transfer between the container and the semi-finished product, air gap, etc.) play an important role during heating and cooling that do not need to be taken into account in conventional heating and cooling without packaging. Therefore, to ensure that the desired temperature of the semi-finished product and a homogeneous temperature field distribution is achieved and the cooling rate during quenching in water is measured, the first semi-finished products were processed with K-type thermocouples attached. The thermocouples were fixed into the container wall (thermocouple 1) and the tool steel blank (thermocouples 2 and 3). The data obtained were then used to design appropriate processing parameters (Figure 2b).

Processing consisted of heating in a furnace at 1250 °C with a holding time of 60 min. The heating was carried out in an electric LAC resistance furnace without a protective atmosphere. Then, the sample was cooled in water for 35 s. This was followed by transfer to a second furnace at 1080 °C with a holding time of 30 min. to equilibrate the temperatures. This heating was included mainly due to the heating of the low-carbon steel container wall, which was very significantly cooled during the quenching in water. A total of 1080 °C is the forging temperature recommended by the material data sheet [25] for X210Cr12 steel. After the material is quenched from the semi-solid state, the microstructure is composed of polyhedral austenite grains and a ledeburitic network. The austenite content is over 90% [10,14,17,18].

Austenite is soft, and the ledeburitic network is very brittle. Therefore, it was necessary to fragment the mesh and distribute it evenly among the austenitic grains. To achieve higher wear resistance, it was further necessary to refine the austenitic grains, transform them into martensite, and promote secondary precipitation of fine carbides. This was performed by forging on a hydraulic press, where the insertion of the deformation in 3, 5, or 10 steps was tested (regimes: 3 × def., 5 × def., 10 × def.), (

Table 2. Processing parameters with transition through semi-solid state.

Regime	Heating Temperature [°C]/Time [min.]	Reheating Temperature [°C]/Time [min.]	Number of Forming Steps	Tempering Temperature [°C]/Time [h]	Cryogenic Treatment Temperature [°C]/Time [h]	Tempering Temperature [°C]/Time [h]
3 × def.	1250/60	1080/30	3	-	-	-
5 × def.			5	-	-	-
10 × def.			10	-	-	-
Temp.			10	300/2	-	-
Cryo			10	-	−160/24	-
Cryo + Temp.			10	-	−160/24	300/2

, Figure 3 and Figure 4a). The reduction size was 55 mm at each step, and the upsetting and elongation steps were alternated so that the semi-finished products were deformed on all sides.

After the deformation was completed, the semi-finished products were quenched in water.

After obtaining very fine microstructures, further possibilities were sought for increasing the fraction of very fine carbides, reducing the fraction of retained austenite, and increasing the toughness of the hard matrix. Two approaches were chosen. In the first one, only tempering at 300 °C for 2 h was used (Regime: Temp.). The temperature of 300 °C was chosen based on information from the material data sheets [26,27] and temperatures reported in the literature [28]. Since an attempt was made to reduce the very high hardness obtained after semi-solid processing, a temperature of 300 °C was used, which resulted in lower hardness values. In the other procedure, the effect of cryogenic treatment at −160 °C with a holding time of 24 h (Regime: Cryo) and the combination of cryogenic treatment and tempering at 300 °C with a holding time of 2 h (Regime: Cryo + Temp.) were tested (Table 2, Figure 4b).

2.2. Methods of Evaluation

The semi-finished products were heated in an electric chamber furnace without a protected atmosphere (LAC s.r.o., Židlochovice, Czech Republic). The semi-finished products were formed on a CKW 6000 hydraulic press with a force of 100 tons. Cryogenic treatment was carried out in a CES-CTC cryogenic box (C.E.S NV, Asse, Belgium), which allows controlled cooling and heating in the temperature range −170 °C to 60 °C. The cooling medium is liquid nitrogen, which is automatically dosed in the required amount by means of a PID controller. The internal dimensions of the cryogenic box are 500 mm × 720 mm × 450 mm, the maximum load capacity of the box is 150 kg.

The processed semi-finished products were cut by a water jet. The microstructures were evaluated by light microscope (SM) and scanning electron microscopy (SEM) on Tescan VEGA 3 microscopes (Tescan, Brno, Czech Republic), and a Zeiss Crossbeam 340-47-44 (Carl Zeiss, Oberkochen, Germany) microscope was also used for more detailed analyses. In addition, hardness measurements were performed according to Vickers HV10 (Wollpert Wilson Instruments, Aachen, Germany). XRD measurements were performed in a Panalytical X’Pert PRO diffractometer (Malvern Panalytical Ltd., Malvern, UK) with Co Kα radiation (lambda = 0.1789 nm) and an X’celerator detector. Measurements were carried out using a divergent beam in the Bragg-Brentano (theta/2theta) vertical geometry. A beta filter was set in the incident beam to improve the signal/background ratio. The measurements were taken with a line focus in the range 30–130 degrees, with a step of 0.05 deg. Phase identification was determined by X’Pert HighScore Plus (Malvern Panalytical Ltd., Malvern, UK) with PDF-4 database and Rietveld refinement by means of Topas V3 (Bruker, Billerica, MA, USA).

Transmission electron microscopy (TEM) on a JEOL 200 CX microscope (JOEL Ltd., Tokyo, Japan) with an accelerating voltage of 200 kV was used for a detailed analysis of the structures. Electron diffraction (ED) was used for phase identification. Single-stage carbon replicas were prepared by grinding on a set of sandpapers and mechanical polishing in diamond suspensions. The microstructure was developed by chemical etching in COR solution. A thin amorphous carbon layer was deposited on the metallographically prepared surface, which was released from the surface by anodic dissolution in 3% Nital at 20 V. The thin foils were thinned from a 1.5 mm thick slice by mechanical grinding on SiC papers to a thickness of about 100 μm. Samples with a diameter of 3 mm were stamped from the slice and prepared to their final thickness by electrodeposition on a TenuPol 5 device (Struers LLC, Cleveland, OH, USA) in 700 mL CH₃OH + 300 mL HNO₃, temperature −20 °C and voltage 15 V.

The WJ2030-2Z-Cobra-PJ60° (water jet cutting machine, from PTV, Hostivice, Czech Republic) was used for the blasting test. The weight loss was evaluated using a KERN EMB 600-2 (Kern and Sohn GmbH, Balingen, Germany) precision laboratory scale with an accuracy of 2 hundredths of a gram. In addition, pin-on-disc tests were performed on equipment from CSM Instruments (CSM Instruments, Peseux, Switzerland).

3. Results and Discussion

3.1. Effect of the Number of Deformation Steps on Grain Refinement

After processing with a transition through the semi-solid state with three deformation steps (Regime: 3 × def.), a microstructure with visible boundaries of the previous austenitic grains with a ledeburitic network was obtained. Although this mesh still bounded the original austenitic grains, it was evident that it had been severely fractured and, in some places, had begun to spheroidize. Very fine secondary carbides were precipitated within the austenitic grains formed by the transition through the semi-solid state. In some grains, the beginning of recrystallization and substructure formation was also detected (Figure 5a). The size of these subgrains was 6.2 ± 3.5 μm. A very high hardness of 748 HV10 was obtained.

With five deformation steps (Regime: 5 × def.), regions with a significantly fine grain size of 4.6 ± 2.1 μm and a very fine carbide precipitation were detected in the microstructure (Figure 5b). However, large areas of the original ledeburitic network were still seen. The hardness value was 633 HV10. X-ray diffraction analysis showed that 60% austenite was still present in the structure.

With ten deformation steps, a very fine microstructure was obtained, which consisted of a martensitic-austenitic matrix with a grain size of 1.4 ± 0.5 μm and fine secondary precipitates (Figure 5c,d). The proportion of austenite decreased to 41%. The hardness value reached up to 738 HV10.

Similar results were found in [19], where deformation was applied in the temperature interval from 1100 to 1000 °C. There was also significant fragmentation of the ledeburitic network, recrystallization, and substructure formation in the original austenitic grains. However, due to the insertion of only three deformation steps, such a significant refinement of the structures was not achieved.

3.2. Tempering

Since a high hardness value of 738 HV10 was obtained after the 10-deformation step regime, which was due to the formation of a large proportion of martensite, it was necessary to carry out tempering to increase the toughness of the material. At the same time, an effort was made to reduce the austenite content, which was still 40% in the structure, and to promote the further formation of very fine chromium carbides.

After tempering at 300 °C/2 h, the matrix showed the character of a tempered structure (Figure 6). The hardness value decreased from 738 HV10 to 663 HV10. The fraction of retained austenite only decreased to 36%.

3.3. Cryogenic Treatment

The primary objective of cryogenic processing was to reduce the austenite content by undercooling below its M_s temperature. Especially for tool steels, cryogenic treatment is accompanied by a positive effect in terms of precipitation of fine dispersed carbides, which cannot be achieved during the standard tempering process. Cryogenic treatment is often used because of the increase in wear resistance. This effect has been published and demonstrated in the literature [20,22,23,29]. Cryogenic treatment was carried out in special furnaces designed specifically for this treatment.

When only cryogenic treatment at −160°C/24 h was applied after the 10 × def. regime, a significant increase in hardness value up to 905 HV10 was observed (Figure 7a). X-ray diffraction analysis confirmed that the austenite content in the structure decreased from the original 41% to 22%. After cryogenic processing, the retained austenite is transformed into martensite, and secondary carbides are precipitated. During cryogenic treatment, the precipitation of secondary carbide (Fe, Cr)₂₃C₆ reduces the chromium and carbon content in the austenite, thereby decreasing its stability. This leads to an increase in the martensite start temperature (M_s) of the austenitic matrix, allowing transformation to martensite. Furthermore, at −196 °C, the retained austenite satisfies the thermodynamic conditions for transformation to martensite [28,30]. All this explains the significant reduction in the fraction of retained austenite after cryogenic treatment.

A further reduction in the austenite fraction was achieved by tempering at 300 °C/2 h. There was also a general tempering of the martensitic structure (Figure 7b), which was reflected by a decrease in the hardness value to 844 HV10. The grain size in both cases was 1.4 μm.

3.4. Transmission Electron Microscopy

Since cryogenic treatment and subsequent tempering results in the formation of additional fine precipitates that play a role in increasing wear resistance, it is necessary to determine their type, quantity, and size. Because these particles are very small, this analysis was carried out on a transmission electron microscope. In addition, a statistical evaluation of the secondary excluded particles was also performed (Equations (1)–(4)).

Individual statistical parameters were evaluated based on the following equations:

$$\mathrm{Medium} \mathrm{particle} \mathrm{size}: D=\frac{total particle length}{number of particles} (\mathrm{m})$$

(1)

$$\mathrm{Number} \mathrm{of} \mathrm{particles} \mathrm{per} \mathrm{unit} \mathrm{area}: N_A=\frac{number of particles}{analyzed area} ({\mathrm{m}}^{-2})$$

(2)

$$\begin{array}{l}\mathrm{Number} \mathrm{of} \mathrm{particles} \mathrm{per} \mathrm{unit} \mathrm{volume}: N_V=\frac{N_A}{t} ({\mathrm{m}}^{-3})\\ (t—\text{replica thickness})\end{array}$$

(3)

$$\mathrm{The} \mathrm{mean} \mathrm{particle} \mathrm{distance} \mathrm{of} \mathrm{the} \mathrm{particles}: L=\frac{1}{\sqrt{N_V·D}} (\mathrm{m})$$

(4)

3.4.1. Structure after Cryogenic Treatment −160 °C/24 h

Analysis using carbon extraction replicas confirmed strong precipitation of particles with different morphologies and sizes (Figure 8). Two types of regions were observed, differing mainly in particle character. These were areas with particles of a regular geometric shape with a high density of precipitates and areas where large particles of irregular shape were precipitated. Due to the size of the particles, there was low extraction into the replicas. In most cases, the presence of particles can be predicted from the indentation left in the replica. Based on statistical analysis, it can be concluded that the particle size range was very wide, ranging from 15 to 4000 nm (Figure 9,

Table 3. Summary of static evaluation of precipitated particles (sample Cryo).

D (10−9 m)	NA (1012 m−2)	NV (1019 m−3)	L (10−7 m)
623 *	2.836 ± 0.645	5.722 ± 0.484	2.515 ± 0.558

* Value is only informative due to the wide range (15 to 4000 nm).

). The histogram of the particle distribution shows a wide interval. One maximum is between 200 and 400 nm and the other between 1200 and 1400 nm.

Based on the solution of the diffraction spectra, the large irregular shaped particles were identified by electron diffraction as M₇C₃ carbide. On the contrary, regular geometric-shaped particles were identified as M₂₃C₆ carbide (Figure 7).

Structural analysis using the thin film method confirmed that the matrix is composed of martensite with very fine needles with an increased density of dislocations or twins (Figure 10a). TEM analysis also confirmed the presence of retained austenite, which was detected between the martensite needles (Figure 10b). Larger irregularly shaped particles were identified as M₇C₃ carbide. In these particles, an increased density of stacking faults was usually observed, which was also reflected in the character of the diffraction points, where deformation in the form of “streaks” was detected due to lattice stresses (Figure 10b and Figure 11). Smaller geometric-shaped particles were identified as M₂₃C₆ carbide (Figure 10b).

3.4.2. Structure after Cryogenic Treatment −160 °C/24 h and Tempering 300 °C/2 h

Next, the sample was analyzed after cryogenic treatment and tempering at 300 °C for 2 h. Structural analysis using carbon extraction replicates showed more significant precipitation than the sample without tempering after cryogen treatment. In addition to large irregular-shaped particles and smaller regular-shaped particles, fine needle-like particles were observed in the matrix, which were mainly precipitated at the boundaries of the martensite laths (Figure 12). The density of the precipitated needle-like particles showed moderate heterogeneity. Large clusters of these particles with different orientations were observed (Figure 12a) or individual needles with precise geometrical orientations (Figure 12b).

The size of large irregularly shaped particles and regular geometrically shaped particles ranged from 15 to 2500 nm (Figure 13a). This interval is related to the particles that were observed only in the TEM, because particles as small as about 7 μm could be detected in the light microscope. The size of the needle-like particles ranged from 9 to 500 nm (Figure 13b). Compared to the sample with cryogenic treatment only, the density of the precipitates is an order of magnitude higher, and the mean distance between particles is an order of magnitude lower (Table 3 and

Table 4. Summary of static evaluation of precipitated particles (sample Cryo + Temp).

D (10−9 m)	NA (1013 m−2)	NV (1020 m−3)	L (10−7 m)
400 *	2.651 ± 0.725	5.302 ± 0.532	0.687 ± 0.152

* Value is only informative due to the wide range (9 to 2500 nm).

). Electron diffraction confirmed the presence of carbides M₂₃C₆ (Figure 14a), M₇C₃ (Figure 14b), and M₃C (Figure 15b).

The thin-film method confirmed that the matrix is composed of tempered martensite with very fine lamellae and an increased density of dislocations or twins (Figure 14b). Larger irregularly shaped particles were identified as M₇C₃ carbide. Fine needle-like particles were observed in the matrix and were identified as M₃C carbide.

3.5. Wear Tests

Materials used in abrasive wear conditions are usually selected according to their microstructure and hardness. However, increased wear resistance can also be achieved by less studied factors. These are grain size, morphology, carbide size, and distribution. Therefore, it is important to subject semi-finished products produced by semi-solid state transition forming technology not only to detailed microstructural evaluation after various processing regimes but also to evaluate them in terms of abrasive wear resistance.

3.5.1. Sandblast Test

The sandblast test was designed specifically for this material under shop conditions and on commonly used equipment. This type of wear could occur when the material is used as a protective plate, e.g., in crusher chambers, tumbling machines, etc.

The sandblast test was carried out on waterjet cutting equipment. A combination of waterjet and abrasive nozzle was used for testing (Figure 16a). The abrasive used was ClassicCut 80, garnet-GMA. The 35 mm × 40 mm specimens were ground to a roughness of 0.8 Ra before abrasion. The test was always performed on two samples (Figure 16b). The weight before and after the test was measured on a laboratory balance. For each sandblast test, 20 g of abrasive was used, and the blasting was carried out for 200 s from a height of 100 mm with a pressure of 4130 MPa.

In addition to the samples processed by the proposed procedures with the transition through the semi-solid state, the experiment was also supplemented with a baseline condition (Figure 16c) and with the condition after conventional heat treatment. The treatment consisted of oil quenching from 960 °C and tempering to 250 °C for 2 h. The processing parameters were selected based on data from the material data sheets [25,26,27]. The structure was composed of tempered martensite and chromium carbides. The hardness was 716 HV10 (Figure 17,

Table 5. Chemical composition of Hardox 450 and Hardox 600 steels (wt.%) [34,35].

Steel	C	Si	Mn	Cr	Ni	Mo	B	P	S
Hardox 450	0.18	0.25	1.3	0.1	0.1	0.04	0.003	0.015	0.001
Hardox 600	0.4	0.56	1.0	1.2	1.5	0.60	-	0.015	0.010

). The sandblast test was also performed on Hardox 450 and Hardox 600 materials, which are currently used in the chambers of crushers and tumbling machines. These are high-quality martensitic steels with high wear resistance [31,32,33], (Table 5). They were tested as delivered with hardness values of 465 HV10 for Hardox 450 and 660 HV10 for Hardox 600 (

Table 6. Results of the sandblast test, including the hardness value.

Sample	Grain Size (μm)	Hardness HV10 (-)	Weight Loss (g)
Initial state	-	216 ± 3	1.73 ± 0.22
Conventional heat treated	-	716 ± 9	1.75 ± 0.05
5 × def.	4.6 ± 2.1	633 ± 3	1.54 ± 0.05
10 × def.	1.4 ± 0.5	738 ± 19	1.40 ± 0.04
Cryo	1.4 ± 0.6	905 ± 5	1.33 ± 0.03
Cryo + Temp.	1.4 ± 0.4	844 ± 14	1.61 ± 0.03
Hardox 450	-	465 ± 8	1.55 ± 0.02
Hardox 600	-	660 ± 9	1.65 ± 0.03

).

The sandblast test showed a significant influence of the nature of the structure and the hardness value. The initial annealed condition with the lowest hardness value of 216 HV10 showed the highest weight loss of 1.73 g of the tested samples (Figure 16c). Even in the case of the conventionally processed sample, there was no reduction in weight loss, which was 1.75 g. Even the presence of large sharp-edged chromium carbides did not lead to better results.

Furthermore, the influence of the number of deformation steps and, therefore, the refinement of the structure was also significant. For the sample with 5 deformation steps, the weight loss was 1.54 g, and for the sample with 10 deformation steps, the measured weight loss decreased to 1.40 g. A further reduction in weight loss was found after applying the cryogenic treatment when the lowest weight loss was found to be 1.33 g. This was mainly due to a significant increase in the hardness value to 905 HV10, the breakdown of retained austenite in the structure, and the promotion of further precipitation of chromium carbides. For the specimens that were tempered after cryogenic treatment, there was an increase in weight loss to 1.61 g, which was due to the tempering of the martensitic structure, and neither further precipitation of chromium carbides nor the presence of another type of carbide led to an increase in sandblast resistance.

Compared to the Hardox 450 and Hardox 600 materials where the weight losses were 1.55 and 1.65 g, respectively, better results were obtained after processing with the transition through the semi-solid state, except for one case with tempering.

3.5.2. Pin-on-Disc Test

The standard pin-on-disc test was also used to determine wear resistance. The tests were carried out at room temperature. A ball with ø 6 mm (Al₂O₃) was used as a pin. Each processing mode was tested two times, and the resulting value is the average of these measurements. Wear area measurements were performed using a profilometer. The result is the average of eight measurements across the wear track. The testing was carried out on a high temperature tribometer from CSM Instruments.

The radius of the track was 1.5 mm, the applied load was 15 N, and the total distance of the track was 141.375 m, which corresponded to 15,000 cycles. The rotation speed was 240 rpm. The testing time was 1.02 h.

The wear rate was calculated after the test according to the following relation [36]:

$$W=\frac{Volume of wear track [{\mu \mathrm{m}}^3]}{Load \left[N\right]·Total distance of track\left[\mathrm{m}\right]}$$

(5)

The same processing modes as the sandblast test were tested, including a comparison with Hardox 450 and Hardox 600 (Figure 18). The results show that the test again replicates the hardness values and is highly dependent on the type of structure tested. The highest through-wear rate of 181.9 × 10⁻⁶ mm³/Nm was found for the initial condition, whose structure consisted of a ferritic matrix, fine cementite, and primary chromium carbides. The lowest values were then achieved after semi-solid processing with 10 deformation steps and after combination with cryogenic processing, where the wear rate ranged from 19.3 to 23.0 × 10⁻⁶ mm³/Nm (Figure 17). The higher wear rate for the specimen after cryogenic processing with a higher hardness value compared to the regime with 10-deformation steps was probably due to the higher brittleness of the material that broke out of the formed traces during the test, which led to the detection of higher wear rate. This breakage is due to the lack of surface plasticity and the more brittle structure compared to the structure without cryogenic treatment. The breakage is generally smaller, but the frequency is higher. In the non-cryogenically processed state, bulkier fragments with a lower frequency can be expected [37]. The Cryo + Temp sample also had a higher tendency to break out due to the higher hardness compared to the sample without cryogenic treatment (10 def.).

Even after this wear test, the standard Hardox 450 and Hardox 600 materials achieved higher wear rates than the proposed semi-solid transition modes. This may be due to the absence of fine precipitates in the martensitic structures of Hardox steels compared to the ledeburitic steel, where there is a large number of fine precipitates of different morphologies in the structure.

4. Conclusions

The X210Cr12 ledeburitic tool steel was processed with a transition through a semi-solid state with forming, which was further combined with cryogenic treatment.

The results show that:

• By inserting 10 deformation steps after the transition to the semi-solid state, a structure with very fine grains of the so-called M-A component with a size of 1 μm and very fine carbides of the M₂₃C and M₇C₃ type can be obtained. The sharp-edged chromium carbides were also removed from the structure.
• Cryogenic processing leads to the precipitation of additional fine carbides and a reduction in the austenite content of the structure from 41% to 22%, resulting in a significant increase in hardness up to 905 HV10.
• The tempering after cryogenic treatment causes the precipitation of M₃C carbide, which had a needle-like morphology compared to the previous carbides.
• The waterjet abrasive sandblast test showed the lowest weight loss after a combination of semi-solid and cryogenic treatment. The samples after this treatment showed up to 40% less weight loss compared to the baseline samples and compared to the Hardox material that is standardly used; the weight loss was 24% less.
• The results were confirmed by the pin-on-disk method. The lowest wear rate was achieved for the samples after combined semi-solid and cryogenic treatment.

The results obtained confirm that the proposed processing method can be used to obtain very fine martensitic structures with fine chromium carbide precipitates in conventionally produced ledeburitic tool steel, which exhibit high wear resistance.