On the Unique Microstructure and Properties of Ultra-High Carbon Bearing Steel Alloyed with Different Aluminum Contents

In this study, ultra-high-carbon steels with 1.4% carbon content alloyed with three different aluminum contents, 2.0%, 4.0% and 6.0%, were studied on their tempering stability and temperature resistance. The results showed that the addition of Al significantly enhanced the tempering stability and temperature resistance of ultra-high-carbon steel. The addition of Al inhibited the transformation of ε-carbide to cementite, suppressed the transition of martensite to ferrite and thus, endowed ultra-high carbon steels to maintain very high hardness during tempering within a wide range of temperature up to 500 °C. The present work provides a useful basis on which to develop bearing steel materials with low density and high hardness.


Introduction
Ultra-high-carbon steels (UHCSs), which are hypereutectoid steels (1 to 2.1% C), have been studied extensively, owing to their unique mechanical properties [1][2][3][4][5][6]. UHCSs have high strength, high hardness and good abrasion resistance but are prone to form network carbides due to carbon segregation, leading to poor plasticity and high brittleness at room temperature [7,8]. However, when prepared by an appropriate process and alloying, UHCSs can be superplastic at elevated temperatures and exhibit high strength and good ductility at room temperature [9,10]. Such intriguing properties are attributed to the elimination of deleterious proeutectoid network carbides and the development of ultrafine carbides in spherical or pearlitic form [11][12][13]. Alloying elements of Al and Si were added to inhibit the precipitation of network carbides in the UHCSs. Due to both Al and Si being graphitized elements, it is necessary to add Cr to stabilize carbides and prevent the precipitation of graphite [14]. At present, the topics on UHCSs are mainly focused on how to obtain an ultrafine grained matrix, how to eliminate network carbides by alloying, superplasticity at elevated temperature and the influence of the spheroidized microstructure on room-temperature mechanical properties. It was revealed that the hardness of UHCS after quenching and tempering is over 65 HRC, which is significantly higher than that of conventional high-carbon bearing steel [4]. In addition, the strength, contact fatigue and wear resistance of the ultra-high-carbon steel are all better than that of the conventional high-carbon chromium-bearing steel [15][16][17][18][19]. The addition of Al significantly reduces the density of steels, due to lattice expansion and the replacement of Fe by Al [20][21][22]. However, at the present time, no detailed study has been demonstrated in the literature on the microstructure evolution during the tempering process, and no further research exists on the effects of aluminum content on the microstructure and properties of ultra-highcarbon steel, the lack of which hinders the research and application of ultra-high-carbon steel, especially in the bearing industry.
When quenched steel is tempered, the ability to resist decreases in strength and hardness is called tempering stability. Temperature resistance refers to the property of maintaining excellent mechanical properties under heated conditions. In order to examine the possibility of the application of ultra-high-carbon steel as a bearing steel, the effect of Al content on tempering stability and temperature resistance of ultra-high-carbon steel were investigated in this study. The addition of Al can effectively inhibit the precipitation of pre-eutectoid network carbides in ultra-high-carbon steel, but excessive addition will cause graphitization to occur, and it is detrimental to hot formability and hot workability. In the existing research, the Al content of ultra-high-carbon steel is generally selected to be 1.6~6%. In order to study the effect of Al on ultra-high-carbon steel, 2%, 4%, and 6% Al is added. In addition, microstructure and hardness were compared and evaluated between the ultra-high-carbon steel and the conventional bearing steel of GCr15. Our aim in this research is to reveal the unique microstructure and properties of ultra-high-carbon-steel alloyed with aluminum and to demonstrate the possibility of the application of ultra-highcarbon steel as a promising bearing material with low density, ultra-high hardness and high temperature resistance.

Experimental
Three types of 1.4% C-UHCSs with different Al content were prepared by induction melting, cast into ingots and forged into bars with a 60 mm diameter. The GCr15 bearing steel for comparison was industrial steel produced by the Xingcheng Special Steel Company (Jiangyin, China).
The chemical compositions of these steels are listed in Table 1. Samples were cut from each bar. The metallographic samples of UHCSs were heated to different temperatures from 700 • C to 1000 • C for 30 min and then quenched with oil to room temperature. The quenched samples were treated at −73 • C for 2 h to reduce the retained austenite content. Finally, the cryogenic samples were tempered at different temperatures from 160 • C to 600 • C for 2 h; a schematic diagram of the heat-treatment process is shown in Figure 1.
The microstructural study was carried out by scanning electron microscopy (SEM, FEI Quanta 650 FEG, FEI, Hillsboro, OR, USA) and transmission electron microscopy (TEM, H-800, HITACHI, Tokio, Japan), and the hardness in each case was obtained using a Rockwell hardness tester (TH-300, TIME, Guangzhou, China). Eight indentations were measured for each sample. In addition, the hardness of the samples tempered at 160 • C was tested after heating for 30 min at 200-500 • C using a high-temperature hardness tester.

Microstructures after Forging and Determination of Phase Transition Temperatures
The microstructures of the UHCSs after forging are lamellar pearlites, as shown in Figure 2. The number and size of the grain boundary carbides decreased significantly with the increase in Al content, which indicated that the addition of Al could obviously inhibit the the network carbides. Because the addition of Al constrained the austenitizing process by suppressing the dissolution of carbides, undissolved carbides are likely to remain after austenitizing. The presence of undissolved carbides must reduce the concentration of solid solution carbon in austenitic matrices. Additionally, the carbon content of austenite decreased with the increase in Al content, which reduced the driving force for the precipitation of carbides and thus, lowered the precipitation of network carbides. In order to design a suitable heat treatment process for UHCSs with different Al content, the phase transition temperature of the materials must be tested first. The phase transition temperature was simply measured by the quenching hardness method. The three types of UHCSs were quenched in oil after isothermal treatment at different temperatures in the range of 750-900 • C for 30 min. Figure 3 shows the SEM microstructure of UHCS-2Al after quenching at different temperatures. The Rockwell hardness-temperature relationship curve of UHCS steel after austenitization + cryogenic treatment + 160 • C × 2 h tempering treatment is shown in Figure 4. The cryogenic treatment is to reduce the content of retained austenite. Among them, Q represents oil-cooled; CT represents cryogenic treatment, and T represents tempered.
Austenizing occurred partly when heated at 800 • C, and the microstructure of quenched UHCS-2Al was a mixture of martensite and pearlite, as shown in Figure 3c. When the heating temperature reached 825 • C, the quenching microstructure was ultrafine martensite, as shown in Figure 3d. The martensite grains grew remarkably when the quenching temperature exceeded 900 • C, and the microstructure is composed of lath martensite and acicular martensite, as shown in Figure 3f.
For UHCS-2Al, the hardness after different quenching temperatures is shown in the black curve in Figure 4a. It is clear that there is a plateau with small fluctuations in hardness value before 775 • C. After that, hardness sharply increases and reaches its peak value at 825 • C. The beginning and end of the sharp increase of hardness both correspond to the start and completion of austenitization, respectively. With further increases in the quenching temperature, hardness gradually decreases, which is due to the coarseness of the martensite grains. Therefore, the Ac 1 and Ac 3 of UHCS-2Al are roughly determined to be 775 • C and 825 • C, respectively. Using the same method, the Ac 1 and Ac 3 of UHCS-4Al are evaluated to be 800 • C and 850 • C, respectively; the Ac 1 and Ac 3 of UHCS-6Al are evaluated to be 825 • C and 875 • C, respectively, as shown in the black curves in Figure 4a.  According to Figure 4c, UHCS-2Al, UHCS-4Al and UHCS-6Al have austenitized at 825 • C, 850 • C and 875 • C, respectively. The peak hardness of the tempered samples with different Al content is between 65-68 HRC. The super-high hardness of UHCSs was beneficial to the wear-resistance and contact-fatigue properties. Examining Figure 4b,c, it is worth noting that the hardness of cryogenic-treated UHCS-2Al steel decreased after tempering, whereas that of UHCS-4Al and UHCS-6Al steels anomalously increased. This indicated that the addition of Al could effectively inhibit the softening of tempering and even increase the hardness of the steel. In addition, the higher the Al content, the more the hardness increases.

Tempering Stability and Temperature Resistance
Based on the above results, three types of UHCSs and GCr15-bearing steel were treated by quenching, cryogenic treatment and tempering at different temperatures to explore the effect of Al content on tempering stability. After the heat treatment, the hardness of the sample was measured, as shown in Figure 5 (Note: tempering at 20 • C indicates the cryogenic treatment).
The hardness of GCr15 steel was 62 HRC after being tempered at 200 • C. When the tempering temperature exceeded 200 • C, the tempering hardness decreased continuously with increasing tempering temperature. The hardness of UHCSs with Al content was stable at higher levels (63-65 HRC) within a certain temperature range. For UHCS-2Al steel, the tempered hardness slowly decreased between 200~360 • C, but the decrease was small, and it was basically maintained at a relatively high level. The hardness of UHCS-4Al steel and UHCS-6Al steel remained basically unchanged with the increase in tempering temperature in the range of 200~400 • C. Only when the tempering temperature exceeded the temperature range did the hardness begin to decrease with increasing tempering temperature. This indicated that the addition of Al inhibited the decomposition of martensite, endowing the UHCSs with good tempering stability. The high-temperature hardness of three types of UHCSs was tested. UHCS-2Al, UHCS-4Al and UHCS-6Al were quenched at 850, 875 and 900 • C, respectively, after cryogenic treatment and tempered at 160 • C for 2 h. The samples were tested using a high-temperature hardness tester at 200, 300, 400 and 500 • C after being heated for 30 min, and the test result is shown in Figure 6, indicating that the three types of UHCSs with Al content had good temperature resistance. The hardness was still maintained at a high level (>58 HRC) when UHCS-2Al steel was heated at 400 • C for 30 min and when UHCS-4Al steel and UHCS-6Al steel were heated at 500 • C for 30 min. This greatly enhanced the potential of ultra-high-carbon steel with Al content as high-temperature bearing steel. The addition of Al greatly improved the tempering stability and temperature resistance of ultra-high-carbon steel, and the higher the content of Al, the more significant the effect was. It was preliminarily considered that the addition of Al greatly inhibited the decomposition of martensite. In order to further verify the above conjecture, the microstructures of three types of UHCSs and GCr15 steels after tempering at different temperatures were observed by SEM. Figures 7-10 show the microstructures of GCr15, UHCS-2Al, UHCS-4Al and UHCS-6Al steel, respectively, after tempering in the range 160-600 • C. Figure 5 shows that the hardness of GCr15 steel decreased relatively slowly when tempering below 360 • C, and the hardness remained at approximately 56 HRC after tempering at 360 • C, indicating that the martensite microstructure was not completely decomposed. Figure 7a did not show obvious granular carbide precipitation on the martensite matrix after the GCr15 steel was tempered at 160 • C. Therefore, there is only a small decrease in hardness after a low-tempering temperature, as shown in Figure 5. Figure 7b-d shows the microstructures of tempering in the range 200-360 • C, indicating that the martensite microstructure did not completely decompose, and some granular carbides precipitated on the martensitic matrix and the microstructure after tempering were tempered martensite and granular carbides, as shown by the yellow arrow in Figure 7. Figure 7e-j shows the microstructures of tempering in the range 400-600 • C. In this temperature range, with the increasing tempering temperature, the tempering martensite decomposed into ferrite and carbides, and the amount of precipitated granular carbides increased and began to spheroidize gradually until completion. With increasing tempering temperature, carbide precipitation increased gradually. The martensitic microstructure decomposed continuously; the carbon content in the matrix continued to decrease, resulting in a continuous decrease in hardness. It was also the main reason why GCr15 bearing steel could not meet the conditions of high-temperature service.   Figure 8a shows that UHCS-2Al steel was tempered martensite after tempering at 160 • C with a small amount of undissolved carbide particles in the quenching process, as shown by the carbide in the green circle in the picture, which are the same as the carbide in Figure 2a. When tempering in the temperature range 200-360 • C, as shown in Figure 8b-e, martensite did not decompose and the number of carbide particles did not obviously increase, thus high hardness was maintained after tempering in this temperature range. When the tempering temperature exceeded 360 • C, martensite began to decompose into ferrite and cementite, and the number of cementite particles increased. When tempering in the range 400-440 • C, as shown in Figure 8f,g, obviously grain boundary carbides appeared, as shown by the yellow arrow, and the grain boundary network disappeared after further increases in the tempering temperature. With the tempering temperature further increased to 520-560 • C, as shown in Figure 8h,i, carbide, as shown by the white arrow, began to spheroidize, and spheroidization was almost complete after tempering at 600 • C, as shown in Figure 8i. As shown in Figure 9, the granular carbides on the martensitic matrix did not increase after UHCS-4Al steel was tempered in the range 200-400 • C, indicating that there were no or only a few cementite particles precipitated. After tempering in this region, the tempering martensite morphology was obviously different from that of GCr15 steel. When UHCS-4Al steel was tempered in the range 200-400 • C, there was a very fine needle carbide precipitate, densely distributed on martensite lath, as shown in the red circle in Figure 9b-e. It was preliminarily considered that the unstable ε−carbides, which were in coherence with the martensitic matrix, precipitated during the tempering process, which will be confirmed later. The ε−carbides were not transformed into cementite particles in this temperature range; therefore, UHCSs with Al content exhibited high hardness after tempering in this temperature range. The microstructure of UHCSs was analyzed in detail by TEM. When tempering temperature exceeded 400 • C, the tempered martensite began to decompose, forming a grain boundary network, as shown by the yellow arrow in Figure 9f,g. With further increases in tempering temperature, granular carbides precipitated and grew, as shown by the white arrow in Figure 9h,j; however, the spheroidizing process for UHCS-4Al steel was not complete until the tempering temperature reached 600 • C. As shown in Figure 10, the microstructure of UHCS-6Al steel after tempering in the temperature range 200-520 • C was similar to that of UHCS-4Al steel after tempering in the temperature range 200-400 • C. During the tempering temperature at 160-440 • C, as shown in Figure 10a-e, the microstructure maintained a martensitic microstructure instead of decomposing into ferrite and cementite. The very fine acicular carbides precipitated along the perpendicular direction to the martensitic lath beam, and no obvious cementite particles precipitated, as shown by the yellow arrow in Figure 10f,g. Due to the conversion of metastable ε-carbide to cementite, the hardness of UHCS-6Al steel was retained above 59 HRC after tempering at 520 • C. The addition of Al inhibited the ε carbides, which were transformed into granular carbides, and very little granular cementite was formed. When the tempering temperature continues to rise, to 560 • C, a small amount of granular cementite can be seen, as shown by the white arrow in Figure 10i.
In essence, the microstructural transformation of high-carbon steel during tempering is divided into several stages. The microstructure of high-carbon steel before tempering was martensite (over the saturated alpha phase) and retained austenite (or partly undissolved carbides). During the tempering process, the temperature range 20-100 • C was the inoculation period. In this temperature region, the C atoms still had some diffusion ability, and the main process was the segregation and aggregation of C atoms. In the range 100-200 • C, the main process was the decomposition of the quenched martensite, i.e., the C atoms from the oversaturated martensite that had dissolved, precipitating the unstable ε-carbides (Fe x C (Fe 2 C, Fe 2.4 C, and Fe 2~3 C), which were in coherence with the martensitic matrix, decreasing the C content, lattice constant C and square degree (c/a ratio) in martensite. In the temperature range 200-350 • C, the main process was the transition of carbide type, namely, the ε-carbides or χ-carbides (Fe 5 C 2 , metastable carbides, which were converted from the ε-carbides at >250 • C) transforming into granular cementites. When the tempering temperature reached higher than 350 • C, the main process was the recovery and recrystallization of the alpha phase and the spheroidization and coarsening of cementites. With increasing tempering temperature, cementites aggregated, grew and gradually converted into spherical particles, and the α-Fe recrystallized and gradually transformed from strip to equiaxed block.
In view of the changes in the microstructures of tempered GCr15 and UHCSs with aluminum content as shown in Figures 7-10, when UHCSs and GCr15 were tempered at 160 • C, the main process was the decomposition of the quenched martensite, i.e., the C atoms from the oversaturated martensite and the unstable nanoscale ε-carbides, which were in coherence with the martensitic matrix precipitated, dissolved and the quenched martensite transformed into tempered martensite. When GCr15 steel was tempered in the range 200-360 • C, the unstable ε-carbides gradually transformed into granular cementites. When the tempering temperature was more than 360 • C for GCr15 steel, with further increasse in tempering temperature, cementites aggregated, grew and gradually converted into spherical particles, and the tempering martensite (alpha phase) further decarbonized and transformed into the ferrite phase, and the ferrite phase recrystallized and gradually transformed from strip to equiaxed block. However, UHCSs with Al content, such as UHCS-2Al, UHCS-4Al, and UHCS-6Al steels, were tempered in the ranges 200-360 • C, 200-400 • C and 200-520 • C, respectively, and there was no obvious conversion of unstable ε-carbides into granular cementites with increasing tempering temperature and prolonged tempering time (no significant increase in granular cementites was observed in the SEM microstructure). Studies [24] have shown that Al increases the eutectoid temperature and carbon content of the eutectoid point of steel, and the increase of the carbon content point of the test steel could make more carbon elements a solid solution in the matrix, thereby inhibiting the segregation of carbides. The addition of a high content of Al prevented the precipitation of cementite and promoted the formation of k carbides. The conversion process of ε-carbide to cementite was suppressed, which improved the tempering stability of ε-carbide. Until the tempering temperature further improved, the cementites precipitated obviously, i.e., ε-carbides began to transform into cementites, and tempering martensite began to transform into ferrite because of further decarbonization. With the further increase in tempering temperature, cementites aggregated, grew, and gradually converted into spherical particles, and the alpha-Fe recrystallized and gradually transformed from strip to equiaxed block.
The addition of Al inhibited or delayed the unstable nanoscale ε-carbides, which were in coherence with the martensitic matrix to transform into cementites, leading to the failure of the martensite to decarbonize to form the ferrite. Therefore, a hardness platform was formed in a certain tempering temperature range; therefore, the UHCSs with Al content exhibited high hardness after a relatively high-temperature tempering and had good temperature resistance. This was the basic reason for the high tempering stability and good temperature resistance of ultra-high-carbon steel due to the addition of aluminum. Moreover, with increasing Al content, the tempering stability and temperature resistance improved. Figures 11-13 show the morphology and diffraction pattern of UHCSs with aluminum content after tempering using a transmission electron microscope (TEM). Through diffraction calibration, it can be determined that the needle-shaped carbides in the picture are ε-carbides (six square structure), as shown in the red circle in Figure 12a-c. Clearly, when UHCS-2Al steel was tempered in the ranges 160-360 • C, as shown in Figure 11a,b, when UHCS-4Al steel was tempered in the ranges 160-400 • C, as shown in Figure 12a-c and when UHCS-6Al steel was tempered in the ranges 160-520 • C as shown in Figure 13a-e, the nanoscale carbides ε-carbides did not decompose, thus fully validating our view. In UHCS-4Al, when the tempering temperature was 480 • C, ε-carbides decreased and turned into other types of carbides, as shown by the yellow circle in Figure 12d; the same microstructure can be seen in Figures 11c and 13f.

Conclusions
The effects of the addition of aluminum on the microstructure and properties of ultrahigh-carbon bearing steel alloys were studied during tempering with temperatures as high as 500 • C. It was revealed that both microstructure and properties were significantly different from the coventional high-carbon bearing steel, such as relative low density, extra-high hardness and high temperature resistance. The main conclusions of this study are as follows.
(1) The addition of Al significantly reduced the density of steel, inhibited the precipitation of carbide networks and significantly improved the phase transition temperature of ultra-high-carbon steel. (2) UHCS-2Al, UHCS-4Al and UHCS-6Al austenitized at 825 • C, 850 • C and 875 • C, respectively, and at low-tempering temperature, the hardness peaks were between 65-68 HRC, and such a high hardness was beneficial for the wear-resistance and fatigue properties. (3) The addition of Al significantly improved the tempering stability and temperature resistance of ultra-high-carbon steel, as the addition of Al inhibited the process of ε-carbide transforming into granular cementite, leading to the failure of martensite to decarbonize to form ferrite. In addition, the greater the amount of added Al, the better the tempering stability and temperature resistance of ultra-high-carbon steel. Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.

Data Availability Statement:
The data presented in this study are available on request from the corresponding author.