Effect of Tempering Temperature on Carbide Evolution and Mechanical Response of Deep Cryogenically Treated Martensitic Stainless Steel

Abstract

Deep cryogenic treatment (DC) is widely applied to martensitic stainless steels to suppress the presence of metastable retained austenite (RA), which may otherwise transform into brittle martensite under deformation and degrade mechanical performance. In this study, a low-carbon 13Cr-2Ni-2Mo martensitic stainless steel was subjected to deep cryogenic treatment for 2 h, followed by tempering at 200–600 °C to investigate carbide evolution and its correlation with mechanical response. At 200 °C, undissolved M₂₃C₆ was observed, accompanied by an RA volume fraction of 8.43% which exhibited a hardness of 543.3 ± 5.1 Hv. When tempered at 400 °C, M₃C became predominant, corresponding to a hardness of 524.5 ± 5.1 Hv. At 500 °C, the simultaneous precipitation of M₃C, M₇C₃, and M₂₃C₆ carbides induced pronounced secondary hardening, which promoted the peak hardness of 559 ± 5.6 Hv. Further tempering at 600 °C resulted in carbide spheroidization M₂₃C₆, which resulted in a hardness reduction to 392.2 ± 3.9 Hv while enhancing ductility. These findings reveal that the tempering temperature plays a decisive role in controlling the carbide precipitation sequence and the stability of retained austenite, thereby enabling the design of an optimal strength–ductility balance in deep cryogenically treated martensitic stainless steels.

1. Introduction

Martensitic stainless steels are utilized in corrosion-resistant plastic molds, dental and surgical instruments, cutlery, and wear-resistant components in the mining industry, owing to their combined corrosion resistance, wear resistance, and mechanical strength [1]. To meet performance demands, microstructure control is critical because retained austenite (RA) is recognized as a soft and metastable phase at low temperatures, which readily transforms into brittle martensite under applied stress or elevated temperature during service. According to the theory of martensitic transformation, this phase transformation is accompanied by a volume expansion which results in dimensional instability that may cause distortion or even locking of bearings during operation [2,3].

Deep cryogenic treatment (DC) has been extensively employed to minimize RA and stabilize the microstructure by promoting the transformation of austenite to martensite [4] and facilitating the precipitation of fine secondary carbides [5,6]. In addition to these established effects, DC has been reported to alter the reversed-austenite formation behavior during subsequent tempering. Zheng et al. [7] demonstrated that for 14.7Cr–6.5Ni–2.0Mo–0.02C super martensitic stainless steel, the RA fraction initially increased and then decreased between 550 and 750 °C, reaching a maximum at 650 °C. The DC-treated steel exhibited a higher and finer RA distribution due to increased Ni-enriched diffusion sites and shortened diffusion paths, while the redissolution of RA at elevated temperatures refined the martensitic laths, thereby enhancing structural stability.

Building on this, recent research on AISI 431 martensitic stainless steel [8] identified a novel cryogenic austenite retransformation (CAR) mechanism, in which the reverted austenite formed during austenite reversion transformation (ART) subsequently retransforms into tertiary α-martensite, ε-martensite, and fine carbides during tempering, leading to enhanced strength and deformation resistance. Moreover, comparative studies on different steels [9] have shown that the effectiveness of DC strongly depends on alloy chemistry and initial matrix structure; the treatment promotes more homogeneous carbide precipitation, particularly when combined with higher austenitizing and lower tempering temperatures.

Therefore, understanding the influence of DC on the microstructural evolution and phase stability of martensitic stainless steels is crucial for optimizing their mechanical and functional properties. In this work, a low-carbon 13Cr–2Ni–2Mo martensitic stainless steel was subjected to deep cryogenic treatment followed by tempering at 200–600 °C to elucidate the carbide evolution sequence and its correlation with mechanical response. The investigation focused on how tempering temperature affects the type, morphology, and distribution of carbides (M₃C, M₇C₃, and M₂₃C₆), as well as the retained austenite stability and hardness variation. The results provide insights into designing an optimal strength–ductility balance in deep cryogenically treated martensitic stainless steels through controlled tempering.

2. Materials and Methods

2.1. Materials Preparation and Heat Treatments

In the present study, 13Cr–2Ni–2Mo martensitic stainless-steel rods were employed as experimental materials. The rods underwent rolling followed by two annealing treatments above the Ac₁ temperature to relieve internal stresses, resulting in a final diameter of 3.5 mm. The chemical composition was verified using Optical Emission Spectroscopy (GVM-1014S, Shimadzu Corporation, Kyoto, Japan) and is summarized in

Table 1. Elemental composition of 13Cr–2Ni–2Mo expressed in weight percent (wt%).

Sample	C	Cr	Ni	Mo	Si	Mn	N	Fe
13Cr–2Ni–2Mo MSS	0.17	12.90	1.80	1.92	0.36	0.08	0.12	Bal.

.

Deep cryogenic treatment (DC) was applied for 2 h in a liquid nitrogen tank after air quenching from 1040 °C. After cryogenic treatment, specimens were tempered in a furnace at temperatures of 200, 400, 500, and 600 °C. The heating rate was maintained at 10 °C/min and each specimen was held isothermally for 2 h. After tempering, all samples were cooled in still air to ambient temperature (~25 °C) without forced convection or quenching. Each batch was processed together to minimize temperature and time variation between samples and are illustrated as shown in Figure 1a. For comparison, conventional quenching and tempering (QT) without DC were also performed under identical austenitizing conditions at 1040 °C for 2 h. The notation used is QCT (tempering temperature) for specimens subjected to DC, and QT (tempering temperature) for samples treated without DC. Dilatometry was employed to determine thermal expansion coefficients and phase transformation temperatures. A DIL 805A/D dilatometer was used with rod specimens (3.5 mm diameter and 10 mm length) under high vacuum (<5 × 10⁻⁴ mbar) to prevent oxidation. The heat treatment simulation involved heating at 10 °C/s to 1040 °C followed by cooling under high-purity helium gas at 100 °C/s to replicate air-cooling conditions. Figure 1b presents the thermal expansion curve of 13Cr–2Ni–2Mo stainless steel. During heating, a slope change at 824 °C marks the onset of austenite formation (AC₁), while a linear trend above 970 °C indicates austenite full transformation (AC₃). Upon cooling from 1040 °C, the rod length shifts from contraction to expansion at 233.91 °C, which corresponds to the martensite start temperature (Ms).

2.2. Microstructure Analysis

Microstructural characterization was performed using field-emission SEM (FEI NOVA NANO SEM 450, FEI Company, Hillsboro, OR, USA) and EBSD (JEOL JSM-7800F Prime FEG-SEM, Tokyo, Japan) with NordlysNano detector (Oxford Instruments NanoAnalysis, High Wycombe, UK). Samples for SEM were sectioned, cold-mounted, and polished with SiC paper (240–4000 grit) followed by alumina suspension (1–0.05 µm). EBSD characterization requires specimens to be electrolytically polished using a Fischione Twin-Jet Electropolisher Model 110 (E.A. Fischione Instruments, Inc., Export, PA, USA) with an electrolyte composed of 15 vol% glycerol, 5 vol% perchloric acid, and 80 vol% anhydrous ethanol. Liquid nitrogen was applied to maintain the electrolyte temperature at −35 °C, while polishing was conducted under a current of 20 A and a voltage of 50 V.

Transmission Electron Microscopy (TEM) is a powerful imaging technique that employs high-energy electron beam transmitted through an ultra-thin specimen, which enables examination of internal structures at a nanometer scale. High-Resolution Transmission Electron Microscopy (HRTEM), an advanced mode of TEM, allows for direct visualization of atomic arrangements, lattice structures, and interfaces. In this study, an FEI Tecnai G2 F20 200 kV FEG-TEM (FEI Company, Hillsboro, OR, USA) was employed to analyze the microstructure. Bright-field imaging, selected-area diffraction patterns, and high-resolution lattice imaging were used to examine the morphology and crystal structures of martensite and precipitates. An integrated Energy-Dispersive Spectroscopy (EDS) system equipped with an Oxford X-MAX 80 detector (Oxford Instruments NanoAnalysis, High Wycombe, UK) was utilized to determine elemental distributions.

TEM specimen preparation involved cutting thin slices approximately 400 μm thick using a precision diamond saw. The slices were mechanically ground using SiC paper of grades 400, 800, and 1000 to a thickness of about 80 μm. Discs with a diameter of 3 mm were then punched to fit the TEM holder and further thinned using finer grades of 1500, 2500, and 4000 to approximately 60 μm. Final thinning was performed by twin-jet electropolishing using a Fischione Model 110 electropolisher (E.A. Fischione Instruments, Inc., Export, PA, USA) under the same electrolyte composition and parameters as the EBSD sample. Polishing was continued until perforation appeared at the center of the disc. The foils were immediately rinsed sequentially with distilled water, 95% ethanol, anhydrous ethanol, and isopropanol to minimize surface contamination and oxidation prior to TEM observation.

Phase identification was further carried out by X-ray diffraction (D2 PHASER, Bruker AXS GmbH, Karlsruhe, Germany) using Cu-Kα radiation (λ = 1.54060 Å, 300 W). XRD scans were conducted in the range of 40–90°, and diffraction patterns were analyzed using MDI Jade6 for phase identification. Retained austenite (RA) volume fraction was calculated from integrated intensities of selected martensite ((211), (200), (110)) and austenite ((111), (200), (220)) peaks using Equation (1) below [10].

$${\mathrm{V}}_{\mathsf{\gamma }}={\displaystyle \frac{1}{1+\left(\frac{{\mathrm{R}}_{\mathsf{\gamma }}^{\mathrm{h}\mathrm{k}\mathrm{l}}}{{\mathrm{R}}_{{\propto }^{\prime }}^{\mathrm{h}\mathrm{k}\mathrm{l}}}\right)\left(\frac{{\mathrm{I}}_{{\propto }^{\prime }}^{\mathrm{h}\mathrm{k}\mathrm{l}}}{{\mathrm{I}}_{\mathsf{\gamma }}^{\mathrm{h}\mathrm{k}\mathrm{l}}}\right)}}$$

(1)

2.3. Mechanical Properties Analysis

Vickers hardness measurements were performed using an FM-810 tester (Future Tech, Kawasaki, Kanagawa, Japan) with a 1 kg load and 10 s dwell time. Samples were embedded along the normal direction, ground with 400–4000 grit SiC, and polished with 1 μm alumina. Five indents were created at 500 μm intervals, and average values were calculated excluding outliers.

Mechanical behavior was evaluated by three-point bending using an MTS 810 UTS (MTS Systems Corporation, Eden Prairie, MN, USA) at 0.48 mm/min. Specimens with a size of 3.5 mm in diameter and 30 mm in length were tested with a 15 mm span and a 4.9 mm indenter. Load–displacement curves were recorded to assess strength, ductility, and fracture behavior.

2.4. Magneto-Induction Method

Magnetic phase (such as martensite and δ-ferrite) content was measured using a ferrite content meter (FERITSCOPE MP30, Helmut Fischer GmbH, Sindelfingen, Germany), which operates based on magnetic induction and eddy current principles. The device allows for non-destructive quantification of ferrite in the range of 0.1–110 FN (0.1–80% Fe) in stainless steels. Measurements were recorded with automatic acceptance at probe contact, where four measurements were taken on each sample to obtain the mean value along with the corresponding standard error.

The fraction of α-martensite and possibly δ-ferrite (f_m) was determined, which used a magnetic field through its probe, thus interacting with the specimen. Since γ-austenite and ε-martensite are paramagnetic [11], only the ferromagnetic α-martensite contributes to the measured response, because the δ-ferrite fraction has no significant contribution to the previous result [12]. Retained austenite measurement used the following method:

$$\mathrm{R}\mathrm{A}\left(\mathrm{\%}\right)=100\mathrm{\%}-\left({\mathrm{F}\mathrm{e}}_{\mathrm{m} }\mathrm{x} \mathrm{k}\right)$$

(2)

where Fe_m denotes ferritescope measurement and k is the curvature correction factor. In the present study, with a 3.5 mm diameter and measured values of 60–80% Fe, a constant correction factor of 1.2 was applied to all samples in present study regardless of the parameter [13,14,15].

3. Results and Discussion

3.1. Microstructure Analysis

After austenitizing at 1040 °C for 2 h, the precipitates present in the as-quenched sample progressively decreased and were eventually dissolved into the matrix, as shown in the SEM micrograph in Figure 2a. In contrast, the conventional quenching and tempering (QT) treatment at 200 °C in sample QT200 preserved finer carbides, since limited atomic diffusion at a low tempering temperature restricted carbide growth and coarsening, as shown in Figure 2b [16,17]. Conventional QT treatment after tempering at 400 °C can be observed in Figure 2c, which reveals needle-type carbide. Tempering at 500 °C and 600 °C showed globular and flaky carbide, as shown in Figure 2d,e, respectively, where at elevated temperatures, enhanced diffusion promoted atom migration to promote more carbides [18].

Although no carbides were detected in the SEM micrograph after quenching at 1040 °C, TEM observations at a higher magnification revealed the presence of very fine nanosized carbides approximately 200 nm in size, which were enriched in chromium and molybdenum, as confirmed by TEM-EDS in Figure 3a and Figure 3b, respectively. The corresponding SAED pattern identified these precipitates as M₂₃C₆ carbides with $\left[1\overline{4}1\right]$ zone axis and a lattice parameter of 1.06587 nm, as shown in Figure 3c. Combining the EDS and SAED results indicates that undissolved M₂₃C₆ (M = Fe, Cr, Mo) persisted in the low-C 13Cr–2Ni–2Mo martensitic stainless steel after quenching.

After tempering at 200 °C for 2 h in sample QCT200, fine globular M₂₃C₆ carbides with a size of 200 nm⁻¹ μm were observed primarily at grain boundaries, as shown in Figure 4a,e. The high dislocation density generated by DC provided preferential nucleation sites, while enhanced carbon diffusion facilitated carbide precipitation and growth compared to the sample without DC, as shown in Figure 2b. At 400 °C, needle-like precipitates (~200 nm) appeared within lath martensite, as shown in Figure 4b,f. Tempering at 500 °C, both needle-like and nanosized globular carbides coexisted, as observed in Figure 4c,g. Tempering at 600 °C, needle-like precipitates diminished while finer globular carbides became more uniformly distributed throughout the matrix, as observed in Figure 4d,h. The superior effect of DC can be ascribed to the enhanced carbide precipitation during tempering at 200–600 °C compared with conventional QT treatment, which was facilitated by the higher dislocation density generated through cryogenic treatment [19,20,21].

As shown in Figure 5, neither the as-quenched condition nor tempering at 200 °C revealed detectable peaks corresponding to carbides or austenite [22]. With tempering at 400–500 °C, martensite peaks became broader due to carbon depletion and fine carbide precipitation, while weak γ(111) reflections appeared, indicating the onset of reversed austenite [23]. At 600 °C, faint M₂₃C₆ peaks emerged, indicating that these carbides reached an XRD-detectable size and fraction. The increase in peak intensity likely reflects a higher volume fraction of M₂₃C₆ [24].

The detailed TEM microstructure of the 13Cr–2Ni–2Mo martensitic stainless steel after 2 h of deep cryogenic treatment followed by tempering at 200 °C was presented in our previous study [12]. In that work, M₂₃C₆ carbides with an average size of about 500 nm (lattice parameter ≈ 1.0677 nm) were identified along the $\left[1\overline{2}1\right]$ zone axis, surrounded by a high dislocation density. These results are consistent with the present SEM observations and provide a useful reference for comparison with the current TEM analysis.

Figure 6a–c presents TEM analysis of 13Cr–2Ni–2Mo MSS tempered at 400 °C, which, as shown in Figure 6a, reveals needle-like precipitates with a size of 170–200 nm in length and 40 nm in width, consistent with SEM observations. SAED performed on these precipitates showed strong and weak diffraction spots arising from two distinct phases. The weaker signals, due to low intensity, were difficult to resolve in dark-field mode, as shown in Figure 6b. The crystallographic alignment between the two phases was identified as $\left[100\right]\mathsf{\theta } // \left[0\overline{1}1\right]\mathsf{\alpha }$, corresponding to the Bagaryatskii orientation relationship commonly observed between cementite (M₃C) and BCC ferrite [25], as illustrated in Figure 6c.

Owing to their nanoscale dimensions after tempering at 500 °C, the precipitates could not be identified using conventional electron diffraction; therefore, high-resolution transmission electron microscopy (HRTEM) combined with Fast Fourier Transform (FFT) analysis was required. The HRTEM bright-field micrograph, as shown in Figure 7a, reveals circular precipitates of approximately 10 nm in size. FFT analysis of the yellow-boxed region, as shown in Figure 7b, confirms martensite with a $\left[\overline{1}13\right]\mathsf{\alpha }$ zone axis, as observed in Figure 7c, whereas the red-boxed region, as shown in Figure 7d, corresponds to the $\left(0002\right)$ and $\left(41\overline{5}1\right)$ planes of M₇C₃ with a $\left[2\overline{3}10\right]{\mathrm{M}}_7{\mathrm{C}}_3$ zone axis, thereby M₇C₃ precipitation within the martensitic matrix after tempering at 500 °C was observed.

In addition to the previously identified carbides, TEM bright-field micrographs, as observed in Figure 7e, and dark-field micrographs, as shown in Figure 7f, revealed the presence of small globular precipitates with a size of 200 nm, which were identified as M₂₃C₆ with a $\left[\overline{1}41\right]{\mathrm{M}}_{23}{\mathrm{C}}_6$ zone axis and lattice parameter of 1.07561 nm. Upon tempering at 500 °C, TEM bright-field micrograph analysis of 13Cr–2Ni–2Mo MSS further revealed needle-like precipitates measuring 130–230 nm in length and 15–25 nm in width, as shown in Figure 7i. The SAED pattern confirmed a Bagaryatskii orientation relationship between ferrite and cementite, consistent with the BCC ferrite diffraction pattern along the $\left[1\overline{3}1\right]$ zone axis as reported by Liu et al. [25,26,27]. Thus, the low-carbon 13Cr–2Ni–2Mo martensitic stainless steel in sample QCT500 contains various types of nanoscale carbides, specifically cementite (M₃C), M₂₃C₆, and M₇C₃.

During tempering at 600 °C, spherical carbides with an average size of approximately 500 nm were detected, as shown in the bright-field and dark-field TEM micrographs in Figure 8a and Figure 8b, respectively. The SAED pattern confirmed the precipitate as M₂₃C₆ with a zone axis of $\left[1\overline{1}1\right]{\mathrm{M}}_{23}{\mathrm{C}}_6$ as observed in Figure 8c, while diffraction pattern analysis, as illustrated in Figure 8d, revealed a lattice parameter of 1.06733 nm for the QCT600 sample. These observations are consistent with the SEM results, which also identified nanosized spherical carbides, as shown in Figure 4g,h, further supported by XRD analysis, where a distinct increase in peak intensity corresponding to the M₂₃C₆ phase was observed. This type of carbide commonly initiates nucleation, as observed in other authors’ findings during high-temperature tempering [28,29,30].

K. Kaneko et al. [31] reported that in 2-Cr martensitic steel, only Fe₃C (M₃C) carbides form within the tempering range of 300–500 °C, which is relatively consistent with the present findings. Tempering at 400 °C, flaky-shaped M₃C (cementite) carbides were observed, whereas at 500 °C, both M₃C and M₇C₃ carbides precipitated simultaneously. In the current study, the formation of M₂₃C₆ carbides was first detected after tempering at 500 °C, which indicates the concurrent presence of M₃C, M₇C₃, and M₂₃C₆ carbides. After further tempering at 600 °C, M₂₃C₆ precipitation became increasingly dominant, which was accompanied by a gradual reduction in the amounts of M₃C and M₇C₃. This transition suggests that the carbide evolution is progressively governed by M₂₃C₆ precipitation, in accordance with the carbide transformation reactions described in the literature [32,33,34,35,36]. Thus, the proposed reaction in the present study is as follows:

Matrix (with undissolved carbide) → M₃C → M₇C₃ → M₂₃C₆

(3)

Figure 9a–f present IPF-Z maps for the as-quenched, QCT200, QCT400, QCT500, and QCT600 samples, respectively. Figure 9f–j show EBSD phase distribution maps of 13Cr–2Ni–2Mo MSS at different tempering temperatures. The quenched sample consists of 0.15% RA, as shown in Figure 9f, and slightly increases after tempering at 200 °C to 0.20%, as shown in Figure 9g. At 400 °C, the RA fraction increases to 0.93% in Figure 9h, which is attributed to carbon partitioning that stabilizes reversed austenite along lath and grain boundaries [37]. However, RA decreases to 0.50% after tempering at 500 °C in Figure 9i as reversed austenite decomposes into ferrite and carbides (mainly M₇C₃) [38]. Further tempering to 600 °C, as observed in Figure 9j, reduces RA to 0.14% due to decomposition into secondary martensite and predominantly M₂₃C₆ carbides [39]. RA fraction measurement from EBSD is summarized in

Table 2. Comparative statistical summary of phase fractions derived from EBSD analysis at various tempering temperatures following 2 h DC.

Tempering Temperatures (°C)	α’ (%)	γ (%)	Carbide (%)
0	98.78	0.15	1.07
200	97.70	0.20	2.10
400	95.72	0.93	3.35
500	96.04	0.50	3.46
600	98.13	0.14	1.73

. Figure 9k–o present the grain boundary misorientation distributions for the as-quenched, QCT200, QCT400, QCT500, and QCT600 samples, respectively. Furthermore, sample QCT500 possesses the highest carbide fraction of 3.46% among all samples, which reflects various types of carbides with relatively high-volume fraction from EBSD measurement.

Figure 10 presents grain boundary misorientation for different tempering temperatures. In this study, a threshold of 10° was adopted to define low-angle grain boundaries, which means that misorientation angles above 10° were excluded from the local misorientation analysis, as they were classified as medium-angle boundaries (10–45°) or high-angle boundaries (>45°) rather than being associated with geometrically necessary dislocations (GNDs) [40]. Tempering at 200 °C in sample QCT200 exhibited the lowest low-angle boundary fraction of 22%, which indicates less recovery and higher stored strain energy during low-temperature tempering [41]. High-angle grain boundaries remained the most prevalent structural feature regardless of tempering temperature, with QCT600 exhibiting the highest fraction of 67.9%, which promotes recovery, reduces dislocation density, and stabilizes high-angle grain boundaries.

A comparison of austenite fractions derived from XRD analysis and EBSD measurements shows a consistent trend across different tempering temperatures, as presented in Figure 11. Both methods confirm an increase in austenite content up to 10.67% at 400 °C for XRD, followed by a decline to 3.36% after tempering at 600 °C. The results indicate that XRD consistently reports higher austenite fractions than EBSD. This discrepancy is likely attributed to the lower spatial resolution of EBSD, which limits its ability to detect fine RA, whereas XRD is a bulk analysis technique that can capture phases more effectively.

3.2. Magnetic Permeability Method

Ferritescope is generally a less sensitive method than other methods like the Vibrating Sample Magnetometer (VSM), since the ferritescope only detects surface area instead of bulk measurement [42]. In the present study, QCT treatment exhibits higher Fe%, hence it possesses a lower RA volume fraction than conventional QT regardless of the tempering temperature applied, as presented in Figure 12. Typically, RA measured by ferritescope has a relatively similar trend with XRD and EBSD results, with slightly higher values compared to XRD. As the tempering temperature increases, the RA fraction decreases in both QT and QCT results. As-quenched samples, QCT200, QCT400, QCT500, and QCT600 exhibit RA fractions of 21.3%, 15.9%, 16.6%, 8.6%, and 6.3%, respectively. In the present study, tempering after 600 °C for sample QCT600 leads to the highest Fe% of 78.1%, resulting in an RA volume fraction of 6.3%, which is 1.8 times higher compared to the XRD result.

The ferritescope consistently reported higher retained austenite values than XRD or EBSD because its magnetic-based measurement is sensitive to all non-ferromagnetic regions, including fine carbides, δ-ferrite, and strain-affected zones within a deep sampling volume (~1–2 mm). These factors locally reduce magnetic permeability and are interpreted as austenite, which results in an apparent overestimation. In contrast, XRD and EBSD quantify only the true FCC phase within limited surface depths, yielding more accurate phase-specific values. Despite these differences, all techniques exhibited consistent trends in RA reduction with cryogenic treatment and tempering, which confirms the reliability of the observed phase transformation behavior [43].

3.3. Mechanical Properties

Figure 13a illustrates the hardness evolution of low-C 13Cr–2Ni–2Mo stainless steel at different tempering temperatures following austenitization at 1040 °C for 2 h. Overall, the QCT process demonstrates a higher hardness compared with conventional QT, primarily due to an increased dislocation density that enhances strengthening [44,45]. In addition, pronounced secondary hardening from more carbide precipitation, as in the SEM result, assists in an increase in hardness [46]. The as-quenched sample exhibits a hardness of 572.3 ± 4.9 HV, which decreases to 543.3 ± 5.1 HV after DC, followed by tempering at 200 °C in sample QCT200. Further tempering at 300°C reduces hardness significantly to 499 ± 4.0 Hv, as carbon diffusion from supersaturated martensite initiates a loss of tetragonality [47]. Tempering at 400 °C in sample QCT400, hardness increases again to 524.5 ± 5.1 Hv owing to cementite (M₃C) precipitation, which contributes to secondary hardening. Maximum hardness, which, in the present study is observed as 559 ± 5.6 Hv, after tempering at 500 °C in sample QCT500, results from pronounced secondary hardening via the precipitation of nanosized M₃C, M₇C₃, and M₂₃C₆ carbides. The observed hardness increases due to dispersion of fine M₂₃C₆ carbides within martensite laths at 500 °C, which promotes secondary hardening [48,49,50]. The secondary hardening observed at 500 °C can be attributed to precipitation strengthening mechanisms, primarily Orowan looping and dislocation pinning [51]. At this temperature, the formation of fine, non-shearable carbides (M₃C, M₇C₃, and M₂₃C₆) effectively impedes dislocation motion, thus increasing the strength of the martensitic matrix [52]. The nanoscale size and uniform dispersion of these carbides, as observed in TEM, are consistent with Orowan-type strengthening, which further explains the hardness peak and subsequent decrease at higher tempering temperatures due to particle coarsening. At 600 °C in sample QCT600, hardness decreases sharply to 392.2 ± 3.9 Hv, which is attributed to excessive grain growth and a consequent reduction in grain boundary strengthening.

Figure 13b presents three-point bending results at various tempering temperatures, with ultimate load and vertical displacement for QT and QCT samples, which are summarized in

Table 3. Comparative table of mechanical properties for QCT and conventional QT samples across different tempering temperatures.

Sample	Ultimate Load (N)	Vertical Displacement (mm)
QCT200	2665.13	11.18
QT200	2772.11	9.15 (fracture)
QCT400	2432.81	18.21
QT400	2720.40	14.47 (fracture)
QCT500	2364.12	4.21 (fracture)
QT500	2370.81	5.99 (fracture)
QCT600	1557.57	18.17
QT600	1574.65	14.87 (fracture)

. Overall, QCT samples exhibited slightly lower maximum loads compared to conventional QT but demonstrated enhanced ductility, with no fracture observed except in sample QCT500. The superior ductility of QCT is attributed to the presence of finer, uniformly distributed carbides and a more stable microstructure which effectively absorbs deformation energy and delays crack propagation [53]. In contrast, QT samples achieved higher peak loads but fractured earlier, which reflects more brittle behavior.

It should be noted that present mechanical evaluation was limited to hardness and three-point bending tests. While these methods reliably indicate relative changes in strength and embrittlement associated with carbide evolution and retained austenite transformation, they do not fully represent the tensile or fracture toughness behavior of the material. Tensile and fracture tests were not conducted due to the limitations of material size and shape, as in the present study, the sample shape is a long rod with a diameter of 3.5mm, where such tests require precisely machined standard specimens (ASTM E8 or E399).

4. Conclusions

This study systematically investigated the effect of tempering temperature on the carbide evolution and mechanical response of 13Cr–2Ni–2Mo martensitic stainless steel subjected to deep cryogenic treatment.

Tempering at 200 °C retained undissolved M₂₃C₆ carbide, while at a higher temperature of 400 °C, cementite (M₃C) began to precipitate. The tempering temperature governed carbide precipitation and retained austenite stability, with fine M₃C, M₇C₃, and M₂₃C₆ carbides forming during tempering at 500 °C. The uniform dispersion of these fine carbides provided strong obstacles to dislocation motion, which resulted in a pronounced secondary hardening effect. At higher tempering temperatures (600 °C), carbide coarsening and partial matrix recovery led to decline in hardness but an improvement in ductility.

Thus, in the present study, tempering at 200 °C in sample QCT200 was the most favorable within the present mechanical scope due to exhibiting a relatively high hardness of 543.3 Hv while maintaining satisfying mechanical properties with relatively lower retained austenite, with values of 8.43% (XRD), 15.9% (ferritescope), and 0.2% (EBSD).