The reliable dissipation of kinetic energy into thermal energy is the fundamental function of a brake disc in high-speed rail systems. This process subjects the disc material to severe thermo-mechanical cyclic loading, leading to complex microstructural evolution directly influencing mechanical properties and service life. Among the materials employed, low-alloy steels processed via casting and forging are predominant. A comprehensive understanding of their phase transformation behavior during non-equilibrium cooling is therefore critical for material selection, processing optimization, and failure analysis. This article presents an in-depth investigation into the continuous cooling transformation (CCT) characteristics of two representative low-alloy steels used for cast and forged brake discs in high-speed trains. The analysis is based on experimental determination of CCT diagrams, microstructural characterization, and hardness evaluation, providing foundational data for predicting microstructural constituents under various cooling conditions akin to service scenarios.
The performance and durability of a brake disc are intrinsically linked to its microstructure. Initial heat treatments typically aim for a tempered martensite or bainitic structure to achieve a favorable balance of strength and toughness. However, during emergency or repeated braking events, the friction surface can experience localized temperatures exceeding the austenitization threshold ($A_{c3}$). Subsequent rapid cooling by ambient air convection can induce non-equilibrium phase transformations, leading to the formation of hard and potentially brittle constituents like martensite or untempered bainite. This localized hardening can create areas of high stress concentration, promoting thermal fatigue crack initiation. Consequently, mapping the transformation kinetics—specifically, the critical cooling rates for the onset of various phases and the resulting hardness—is essential. The science of steel casting allows for the production of complex, near-net-shape components like brake discs, but the solidification structure and inherent segregation can influence its transformation behavior compared to wrought, forged counterparts. This study delves into these comparative aspects.
1. Material Composition and Experimental Methodology
The investigation focuses on two commercially used low-alloy steels, one designated for steel casting brake disc applications and the other for forged brake discs. Their chemical compositions, obtained via optical emission spectrometry, are detailed in Table 1. The compositional differences, though subtle, are strategically designed to achieve specific hardenability and high-temperature strength profiles suitable for their respective manufacturing routes and service demands.
| Element | Cast Brake Disc Steel (wt.%) | Forged Brake Disc Steel (wt.%) |
|---|---|---|
| C | 0.26 | 0.30 |
| Si | 0.58 | 0.40 |
| Mn | 1.04 | 0.75 |
| Cr | 0.90 | 1.20 |
| Ni | 1.00 | 0.40 |
| Mo | 0.55 | 0.65 |
| V | 0.10 | 0.30 |
| P | 0.014 | 0.009 |
| S | 0.011 | 0.007 |
The critical phase transformation temperatures, namely the austenite finish temperature ($A_{c3}$) and the martensite start temperature ($M_s$), were determined using a high-resolution dilatometer (DIL 805L). Cylindrical specimens with dimensions of Ø4 mm × 10 mm were subjected to controlled heating and cooling cycles. For the determination of CCT diagrams, specimens were austenitized at 900°C for 10 minutes to ensure a homogeneous single-phase austenitic microstructure and to dissolve carbonitride precipitates. Subsequently, they were cooled to room temperature at precisely controlled linear cooling rates ranging from 0.2 °C/s to 100 °C/s. The corresponding dimensional change (dilatation) was continuously recorded. Phase transformation start and finish points were identified using the tangential method on the dilation-temperature curves.
Post-dilatometry, the specimens were sectioned, polished, and etched for metallographic examination using optical microscopy (OM) and scanning electron microscopy (SEM). Microhardness measurements were performed using a Vickers indenter under a standard load to correlate the microstructure with mechanical properties. The combined data from dilatometry, microscopy, and hardness testing were used to construct the comprehensive CCT diagrams.
2. Critical Temperatures and CCT Diagram Construction
The analysis of dilation curves during heating and very fast cooling (100 °C/s) provided the fundamental critical temperatures. For the low-alloy steel designated for steel casting, the austenitization completion temperature $A_{c3}$ was found to be 847°C, and the martensite start temperature $M_s$ was 347°C. In contrast, the forged steel exhibited a higher $A_{c3}$ of 862°C and a higher $M_s$ of 368°C.
The displacement of these critical points can be rationalized through the influence of alloying elements on the thermodynamics of iron-carbon alloys. The $A_{c3}$ temperature is primarily governed by elements that alter the stability ranges of ferrite and austenite. Nickel and manganese, present in higher amounts in the steel casting alloy, are austenite stabilizers. They lower the $A_{c3}$ temperature by expanding the austenite phase field in the Fe-C phase diagram. Conversely, strong carbide-forming elements like vanadium and molybdenum, present in higher concentrations in the forged steel, stabilize ferrite and raise the $A_{c3}$ temperature. The $M_s$ temperature is predominantly controlled by the carbon content and the strength of the austenite prior to transformation. Carbon strongly suppresses $M_s$. Although the forged steel has a slightly higher carbon content (0.30% vs. 0.26%), the significant difference in $M_s$ (21°C higher) suggests a more pronounced effect from other solutes. The higher nickel in the cast steel significantly increases the hardenability by stabilizing austenite and depressing the $M_s$ temperature to a greater extent. The $M_s$ temperature can be empirically estimated using equations like the following, which account for alloying effects:
$$
M_s(°C) = 539 – 423C – 30.4Mn – 17.7Ni – 12.1Cr – 7.5Mo + 10Co – 7.5Si
$$
Plugging in approximate average compositions highlights the strong depression effect of Ni and C. The final constructed CCT diagrams for both steels, incorporating all transformation start and finish lines, are the central results of this experimental work.
3. Microstructural Evolution and Hardness under Continuous Cooling
The microstructural constituents observed after continuous cooling at various rates are systematically summarized below, along with their corresponding Vickers hardness (HV) values.
For the Cast Brake Disc Steel:
The CCT diagram reveals a notably suppressed polygonal ferrite and pearlite transformation region. Even at the slowest cooling rate of 0.2 °C/s, the microstructure consists primarily of granular bainite (GB), characterized by irregular ferrite grains with dispersed M-A (Martensite-Austenite) constituents. This indicates the high hardenability imparted by the alloy design for steel casting. As the cooling rate increases to 1-2 °C/s, a mixture of granular bainite and lath-like bainite (LB) forms. In the intermediate cooling rate range of 5-20 °C/s, the transformation products shift to a mixture of lath bainite and martensite (M). The volume fraction of martensite increases progressively with cooling rate. The critical cooling rate for achieving a fully martensitic microstructure in this steel casting material is approximately 20 °C/s. Beyond this rate, the microstructure is predominantly lath martensite, and further increases in cooling rate (up to 100 °C/s) do not alter the microstructural identity, only refining the martensite lath packet size.
For the Forged Brake Disc Steel:
The forged steel exhibits a small but distinct ferrite-pearlite (F+P) nose in its CCT diagram. At 0.2 °C/s, a predominantly ferrite-pearlite microstructure with minor bainitic areas is observed. This contrasts sharply with the cast steel behavior at the same rate. In the cooling rate range of 1.0 to 20.0 °C/s, the microstructure is a complex mixture of bainite and martensite. The critical cooling rate for obtaining a 100% martensitic structure is significantly higher, at about 50 °C/s, compared to the steel casting alloy.
The hardness evolution as a function of cooling rate and microstructure is a key practical output. The trends for both materials are captured in Table 2 and the subsequent analysis.
| Cooling Rate (°C/s) | Cast Steel Microstructure | Cast Steel Hardness (HV) | Forged Steel Microstructure | Forged Steel Hardness (HV) |
|---|---|---|---|---|
| 0.2 | Granular Bainite (GB) | ~383 | Ferrite + Pearlite + (Minor Bainite) | ~220 |
| 1.0 | GB + Lath Bainite (LB) | ~490 | Bainite + Martensite | ~450 |
| 5.0 | LB + Martensite (M) | ~505 | Bainite + Martensite | ~480 |
| 20.0 | Predominantly M | ~520 | Bainite + Martensite | ~510 |
| 50.0 | Fully Martensite | ~525 | Fully Martensite | ~540 |
| 100.0 | Fully Martensite | ~529 | Fully Martensite | ~545 |
For the steel casting material, hardness increases dramatically from 383 HV to 490 HV as the cooling rate rises from 0.2 to 1.0 °C/s, corresponding to the transition from granular bainite to a harder mixture containing lath bainite. The subsequent increase in hardness from 490 HV to about 520 HV in the 1.0-20.0 °C/s range is more gradual, attributable to the increasing martensite fraction. Once the fully martensitic structure is achieved at 20 °C/s, hardness saturates, with only marginal increases due to slight refinement. The forged steel shows a very low hardness at 0.2 °C/s due to the soft ferrite-pearlite structure, but its hardness rises steeply upon entering the bainite-martensite formation regime, eventually surpassing the hardness of the cast steel at very high cooling rates due to its higher carbon content in martensite.
4. Discussion on Transformation Mechanisms and Alloying Effects
The stark difference in the position of the ferrite-pearlite “nose” between the two steels is a direct consequence of alloy design. Hardenability, which is the resistance to the formation of soft diffusion-controlled phases like ferrite, is crucial for steel casting components that may experience moderate cooling rates after casting or during heat treatment. The cast steel’s composition is rich in hardenability enhancers: Manganese (1.04%) and Nickel (1.00%). Nickel, being a non-carbide former, segregates to the austenite, increasing its stability and significantly retarding the nucleation and growth of proeutectoid ferrite. Molybdenum (0.55%) further suppresses diffusion-controlled transformations by segregating to austenite grain boundaries. Although vanadium is a potent hardenability agent when in solution, at 0.10% it primarily forms fine carbonitrides that pin grain boundaries but have a lesser effect on moving the ferrite nose to longer times than solutes like Ni and Mo.
In contrast, the forged steel relies more on chromium (1.20%) and vanadium (0.30%) for hardenability and secondary hardening. While Cr increases hardenability, its effect on suppressing the ferrite nose is less potent than Ni. The higher V content leads to more pronounced precipitation, which can paradoxically provide nucleation sites for ferrite if not fully dissolved during austenitization, potentially explaining the presence of the small F+P nose. The higher $A_{c3}$ of the forged steel is consistent with its lower overall austenite-stabilizer content (lower Mn, much lower Ni) and higher ferrite-stabilizer content (higher Cr, V).
The transformation kinetics can be conceptually described by the Johnson-Mehl-Avrami-Kolmogorov (JMAK) equation for isothermal transformation, but adapted for continuous cooling. The start time $t_s$ for a phase transformation under continuous cooling can be related to the cooling rate $\dot{T}$ and the isothermal transformation start time from a TTT diagram, often expressed as:
$$
\int_{0}^{t_s} \frac{dt}{\tau(T(t))} \approx 1
$$
where $\tau(T)$ is the characteristic incubation time at temperature $T$. A higher hardenability, as seen in the steel casting alloy, means $\tau(T)$ for ferrite is very large, so the integral never reaches 1 before the temperature falls below the nose, bypassing the transformation entirely.
The lower $M_s$ of the cast steel is a critical finding. A lower $M_s$ means martensite forms at a lower temperature, where the yield strength of the austenite is higher, leading to a higher transformation stress and a greater propensity for auto-tempering or even quench cracking. This must be considered in the heat treatment design for such steel casting components.
5. Practical Implications for Brake Disc Performance and Life
The CCT diagrams serve as a predictive map for microstructural changes during in-service thermal cycling. During a severe braking event, the surface layer of the brake disc can heat rapidly above the $A_{c3}$ temperature, creating a superficial layer of austenite. The subsequent cooling rate, driven by conduction into the bulk disc and convection to air, is not uniform. It can range from very slow in the bulk to moderately fast (potentially tens of °C/s) at the very surface.
For the forged steel disc, if a surface region cools at a rate between ~1 and 50 °C/s, it will transform into a mixture of bainite and martensite, leading to a localized hardness increase (e.g., from a base hardness of ~300 HV to over 500 HV). This creates a hard, brittle layer on the surface. For the steel casting disc, the transformation under similar cooling conditions would also lead to bainite and martensite, but the base hardenability is already higher. The critical insight is that the steel casting alloy requires a lower cooling rate (20 °C/s vs. 50 °C/s) to form a fully martensitic, potentially more brittle, surface layer.
From a material selection perspective for brake discs, a higher $A_{c3}$ temperature and a higher critical cooling rate for martensite formation are generally desirable. A higher $A_{c3}$ means a higher surface temperature is required to austenitize the material, making microstructural transformation less likely during normal braking cycles. A higher critical martensite cooling rate makes it less probable that the post-braking cooling will result in 100% untempered martensite, promoting instead a tougher bainitic or mixed microstructure. According to this criterion, the forged steel in this study has a marginal advantage ($A_{c3}$: 862°C vs. 847°C; Critical rate for M: 50°C/s vs. 20°C/s). However, the steel casting process offers immense geometric and economic advantages for complex shapes like ventilated brake discs. Therefore, the alloy design for steel casting must meticulously balance hardenability for uniform properties after the casting process with transformation stability under service conditions to minimize damaging phase changes.
The localized hard zones resulting from phase transformation have a different coefficient of thermal expansion and lower fracture toughness compared to the base material. This mismatch promotes thermal stress concentration and facilitates the initiation of thermal fatigue cracks, often observed as network (“craze”) cracking on brake disc surfaces. Understanding the CCT behavior allows engineers to model and predict the depth and severity of this transformed layer, informing maintenance schedules and disc life predictions.
6. Conclusion
The continuous cooling transformation behavior of two low-alloy steels for high-speed train brake discs has been systematically characterized. The steel designated for steel casting applications exhibits a higher hardenability, evidenced by a suppressed ferrite-pearlite transformation region and a lower critical cooling rate (20 °C/s) for achieving a fully martensitic microstructure. Its $A_{c3}$ and $M_s$ temperatures are 847°C and 347°C, respectively. The forged steel counterpart shows a discernible ferrite-pearlite nose, a higher $A_{c3}$ (862°C), a higher $M_s$ (368°C), and requires a faster cooling rate (50 °C/s) for complete martensite formation.
These differences are governed by the distinct alloying strategies: the steel casting alloy leverages higher Ni and Mn for austenite stability and hardenability, while the forged steel employs higher Cr and V for strength and secondary hardening. The derived CCT diagrams provide an essential tool for predicting in-service microstructural evolution. During braking, localized austenitization followed by rapid cooling can lead to the formation of hard bainitic and martensitic phases on the friction surface, increasing hardness and susceptibility to thermal fatigue cracking. The results underscore that optimal brake disc material design, particularly for the versatile steel casting route, must carefully consider not only static mechanical properties but also the phase transformation kinetics under realistic thermal cycles to ensure long-term structural integrity and reliability.