Optimization of Heat Treatment Parameters for Austenitic Stainless Steel Cladding Using the Taguchi Method

Abstract

Hot-rolled A283 Gr C carbon steel/A240 TP 316L stainless steel-clad plates are widely used in structural applications. However, the hot-rolling process introduces residual stresses and microstructural heterogeneities near the interface, which can adversely affect mechanical performance. This study aims to optimize stress-relief annealing parameters for hot-rolled A283 Gr C/A240 TP 316L-clad steel in order to enhance toughness while preserving microstructural integrity. A Taguchi experimental design based on an L9 orthogonal array was employed to evaluate the effects of holding temperature, holding time, and heating/cooling velocity on Charpy impact toughness. Signal-to-noise (S/N) ratio analysis and ANOVA were used to identify the most influential parameters. Microstructural observations, microhardness profiling, and Charpy impact testing were conducted before and after heat treatment. The results indicate that stress-relief annealing does not alter the base microstructures of either the carbon steel substrate or the austenitic stainless steel-clad layer, nor does it induce carbide precipitation or secondary phase formation in the A240 TP 316L stainless steel. A noticeable reduction in the thickness of the decarburized ferrite zone near the interface was observed, suggesting improved interfacial stability. Microhardness measurements revealed a moderate decrease in hardness near the interface, accompanied by a significant increase in Charpy impact toughness under optimized conditions. ANOVA results show that holding temperature is the dominant factor influencing toughness, followed by heating/cooling velocity, while holding time has a minor effect. The optimal stress-relief annealing conditions were identified as 550 °C for 45 min, with a heating/cooling velocity of 100 °C/h. These findings demonstrate that the Taguchi method is an effective approach for optimizing heat treatment parameters and improving the mechanical integrity of hot-rolled stainless steel-clad plates.

1. Introduction

The demand for high-performance, cost effective materials has driven the universal adoption of clad steels in industrial applications. These composites consist of a low-carbon steel base (e.g., A283 Gr C) metallurgically bonded to a corrosion resistant layer (e.g., A240 TP 316L), combining mechanical strength with enhanced durability [1,2,3,4]. Clad steels offer significant economic advantages over monolithic stainless steel, particularly for thick section components like chemical tanks and heat exchangers [5,6]. Their production involves cladding techniques such as hot/cold-roll bonding, weld overlay, or explosive welding, which ensure metallurgical compatibility but may introduce microstructural heterogeneities at the bonding interface [2,4,7]. These localized effects can influence the material’s thermomechanical behavior, which is a critical consideration for high temperature service applications [8].

Heat treatments are key to enhancing the mechanical properties and microstructural stability of steel components, especially after fabrication. Through controlled heating and cooling, these treatments modify the steel’s microstructure by refining grains, inducing phase transformations, and optimizing precipitation to achieve superior strength, ductility, and toughness [9,10]. For clad steels like A283 Gr C (base) and A240 TP 316L (cladding), thermal exposure modifies the bonding interface via elemental diffusion, microstructural changes, and mechanical cohesion. These are critical factors for composite performance. Optimizing heat treatments requires balancing the distinct behaviors of:

A283 Gr C: High thermal expansion, ferritic–pearlitic structure;

A240 TP 316L: Austenitic stability, lower expansion, and Cr/Ni diffusion kinetics.

Mismatched thermal cycles can lead to interfacial degradation, residual stress accumulation, or even delamination in clad systems. These phenomena are often intensified by differential phase transformations and diffusion driven heterogeneities at the bonding zone [7,8]. Therefore, understanding and evaluating the thermal response of each constituent material individually—particularly in terms of phase stability, transformation kinetics, and interfacial interactions—is essential to ensure interfacial integrity and overall mechanical reliability.

Austenitic stainless steel A240 TP 316L exhibits excellent mechanical strength and corrosion resistance. However, its performance is highly sensitive to thermal history, requiring careful control during heat treatment to avoid deleterious phase transformations [11].

Annealing at 1050–1100 °C is an example of the heat treatments applied on A240 TP 316L steel. This treatment can effectively dissolve the chromium carbides (M₂₃C₆) formed after welding or cold working and restores corrosion resistance by preventing sensitization. However, prolonged exposure at these temperatures may promote excessive grain growth and reduce ductility and toughness [11,12].

The temperature range of annealing between 500 and 900 °C is particularly critical to the austenitic stainless steel. Temperatures approaching 650 °C can initiate chromium carbide precipitation along grain boundaries and promote σ-phase formation, which is a brittle Fe–Cr–Mo intermetallic phase. These mechanisms lead to sensitization, intergranular corrosion, and embrittlement [8,12,13].

In addition, stress-relief treatments of the austenitic steel A240 TP 316L in the range of 425–550 °C can effectively reduce residual stresses while limiting carbide precipitation. However, prolonged exposure above approximately 485 °C should be avoided as it may enhance surface oxidation and delta-ferrite decomposition. These effects can further influence interfacial stability in clad steel configurations [14,15].

In contrast to austenitic stainless steels, low-carbon steel A283 Gr C exhibits a simpler thermal behavior dominated by ferrite–pearlite transformations. Nevertheless, appropriate heat treatment remains essential to optimize mechanical properties and to minimize interfacial mismatches in clad systems [16,17].

Due to its low carbon content (≤0.24%) and limited hardenability, A283 Gr C responds favorably to conventional heat treatments. Normalizing at 850–950 °C, followed by air cooling, refines the microstructure through austenite-to-ferrite/pearlite transformation. This process improves the strength–toughness balance by reducing grain size and homogenizing phase distribution [7,16].

The full annealing at 800–850 °C of low-carbon steel promotes optimal ductility and machinability. This treatment induces controlled ferrite–pearlite coarsening and cementite spheroidization, making it suitable for pre-forming and welding operations while preserving phase stability [13].

Post-fabrication, stress-relief annealing is particularly effective for low-carbon steels in the temperature range of 540–620 °C. In this range, residual stresses are reduced through recovery mechanisms without altering the ferritic–pearlitic matrix [17]. Careful temperature control is required, as subcritical annealing below 650 °C preserves grain stability and prevents unwanted austenite reformation. However, exceeding this threshold may degrade mechanical properties due to excessive grain growth or partial austenitization [18]. Then, unlike austenitic stainless steels, A283 Gr C does not undergo intermetallic phase formation.

While carbon and austenitic stainless steels exhibit distinct thermal responses when processed independently, their combination in clad systems introduces interfacial complexities requiring specialized heat treatment protocols. In clad configurations, heat treatment can promote carbon diffusion across the bonding interface. This diffusion may lead to localized decarburization in the carbon steel and carburization in the 316L layer. These compositional gradients modify the local phase distribution. In some cases, they promote the formation of brittle martensitic or pearlitic zones. Such heterogeneities can reduce interfacial cohesion and decrease fracture resistance [14]. The bonding interface between the stainless steel cladding and the carbon steel substrate is then a critical region. In this zone, complex metallurgical interactions occur during post-cladding heat treatments. Elemental diffusion plays a dominant role in the interfacial microstructural evolution. This diffusion mainly involves carbon, chromium, nickel, and iron, and directly affects the mechanical properties and long-term performance of the clad material.

In previous studies, Jiang et al. [19] investigated the heat treatment of stainless steel-clad plates within a temperature range of 500–700 °C. This treatment promoted interfacial diffusion, leading to the formation of a martensitic layer at the bonding interface. This layer exhibits significantly higher hardness than both the carbon steel substrate and the stainless steel cladding [19]. The thickness of the interfacial diffusion zone increases with prolonged holding times. It also grows under repeated thermal cycles, further influencing local mechanical heterogeneity. Such microstructural modifications frequently culminate in localized embrittlement and diminished fracture toughness at the interface, presenting a critical reliability concern for clad structures.

Then, a great effort is required in the choice and optimization of heat treatment conditions. For example, thermal processing at 620 °C for 1 h has demonstrated particular effectiveness, achieving up to a 50% reduction in residual stress while minimizing adverse microstructural changes, as reported in [20]. Comparable stress-relief annealing can be attained at lower temperatures such as 540 °C, but it requires prolonged dwell times. This approach offers the added advantage of restricting grain growth and preserving interfacial toughness. The optimization of heat treatment parameters proves equally essential in balancing the stainless steel cladding’s corrosion resistance against the carbon steel substrate’s mechanical strength. For instance, combined quenching and tempering sequences, such as high temperature austenitization at ~1080 °C followed by tempering at 550 °C, can effectively reconcile these properties from the perspective of inducing pronounced phase transformations in the base steel near the interface [21].

It is important to emphasize that interfacial toughness is strongly influenced by the applied thermal history, as well as by the metallurgical characteristics of the bonding zone. Notably, clad systems incorporating nickel interlayers exhibit distinct crack propagation mechanisms and a higher susceptibility to hydrogen embrittlement compared with nickel-free configurations, which generally demonstrate superior toughness under ambient conditions [22].

Accordingly, the optimization of heat treatment parameters plays a pivotal role in enhancing the mechanical performance of steels. In this context, the Taguchi method has been widely employed as an efficient statistical optimization approach, frequently coupled with analysis of variance (ANOVA) to identify significant factors and establish optimal processing conditions. For instance, Agboola et al. [23] investigated the optimization of heat treatment parameters for medium-carbon steel using an L9 orthogonal array coupled with ANOVA. Their results revealed that the quenching medium, soaking temperature, and soaking time exert significant effects on mechanical properties, including hardness, yield strength, and ultimate tensile strength [23].

In clad systems, the interaction between interfacial phenomena and thermal parameters is complex. This complexity requires an efficient and systematic optimization approach. The Taguchi method offers a statistically robust framework to determine optimal heat treatment conditions while reducing the number of experimental trials [24]. By using orthogonal arrays, the Taguchi method enables structured evaluation of key parameters, including temperature, time, and cooling velocity, as well as their interactions [25]. This makes the method particularly suitable for clad steels optimization, where multiple performance criteria, such as interface integrity and mechanical properties, must be balanced.

In this context, Karna and Sahai [26] provided a comprehensive overview of the Taguchi method, highlighting the use of orthogonal arrays for efficient experimental design with minimal trials [26]. They emphasized the signal-to-noise (S/N) ratio as a key metric for robust optimization, allowing for the maximization of desired responses while minimizing variability caused by noise factors.

In Karimi et al. [27], a Taguchi L16 orthogonal array was applied to optimize the deposition of Ni–WC MMC overlays on steel substrates. Wire feed speed, travel speed, and shielding gas conditions were identified as the most influential parameters. These factors significantly affected carbide retention, dilution, deposition rate, and bead geometry. The results showed that these responses were mainly governed by process parameters rather than heat input alone. Overall, the study demonstrates that the Taguchi approach is an efficient and systematic tool for improving wear-resistant cladding performance [27].

In the same context, Zhu et al. [28] applied the Taguchi method. This systematic methodology provides a robust framework for optimizing the laser cladding parameters of 316L stainless steel and enhancing the overall quality of the deposited coatings. Laser cladding is a widely used surface engineering technique for improving the mechanical and corrosion resistance of 316L stainless steel through localized heat input and rapid solidification [28].

With the aim of optimizing heat treatment parameters, Zalilah et al. [29] have employed the Taguchi method using an L9 orthogonal array. The experiments evaluated temperature (400–800 °C), heating time (60–120 min), heating rate (10–20 °C/min), and substrate roughness (P80, P240, P800). Optimal conditions were identified as 800 °C, 120 min, 10 °C/min, and P80 roughness, which maximized both thickness and hardness while enhancing surface quality to prevent delamination and corrosion in carbon steel fiber metal laminate applications [29].

The present study investigates the effects of stress-relief annealing on the mechanical and metallurgical properties of hot-rolled A283 Gr C/A240 TP 316L-clad steel plates. Employing a Taguchi L9 orthogonal array design, this research systematically evaluates the influence of three key parameters—holding temperature, holding time, and heating/cooling velocity—on microstructural evolution, microhardness profiles, and Charpy impact toughness. The primary objective is to identify optimal annealing conditions that maximize toughness and moderate microhardness near the interface, while preserving microstructural integrity and preventing interfacial degradation. Signal-to-noise ratio analysis and ANOVA are utilized to quantify parameter contributions and interactions, ensuring a statistically robust optimization. The following sections outline the experimental methodology, present the results, and analyze the underlying mechanisms.

2. Materials and Methods

2.1. Austenitic Stainless Clad Steel Material

The hot-rolled A283 Gr C/A240 TP 316L-clad steel plates were fabricated through hot-roll bonding, combining a 12 mm thick low-carbon steel substrate (A 283 Gr C) with a 3 mm thin austenitic stainless steel cladding layer (A 240 TP 316L). The geometrical characteristics of the investigated austenitic stainless clad steel specimens are illustrated in Figure 1. The chemical composition is presented in Table 1.

Figure 1. Photograph of the hot-rolled clad steel plate showing the A283 Gr C carbon steel substrate (12 mm thickness) and the A240 TP 316L stainless steel-clad layer (3 mm thickness).

Table 1. Chemical composition of the A283 Gr C carbon steel substrate and the A240 TP 316L stainless steel-clad layer (wt.%).

Elements	C	Si	Mn	P	S	Cr	Ni	Mo
A283 Gr C	0.167	0.172	0.479	0.019	0.022	-	-	-
A 240 TP 316L	0.017	0.514	1.124	0.038	0.002	16.654	10.109	2.092

2.2. Mechanical and Microstructural Characterization of the Austenitic Stainless-Clad Steel

2.2.1. Microstructural Characterization

Specimens with dimensions 30 × 30 × 12 mm³ were used to perform the microstructural examinations of the substrate, the clad layer, and the bi-material interface. Metallographic sample preparation was performed using a GP-1B metallographic grinding and polishing machine (Laizhou Weiyi Experiment Machine Manufacturing Co., Ltd., Laizhou, China). The specimens were mechanically polished using SiC abrasive papers with grit numbers varying from 280 to 1200, followed by a final polishing using 6, 3, and 1 µm diamond pastes, respectively, as depicted in Figure 2a. Subsequently, these specimens were chemically etched at room temperature using nital 3% for 10 s for the carbon steel side, while the stainless steel side was etched using 75% HCL/25% HNO₃ (aqua regia) for 6 s. Metallographic examinations were conducted before and after the application of heat treatments in accordance with ASTM E112 [30] and ASTM E3 [31]. Microstructural observations were carried out using a Leica DM ILM metallographic microscope (Leica Microsystems GmbH, Wetzlar, Germany) operating in reflected light mode (Figure 2b).

Figure 2. Equipment used for metallographic observation and mechanical testing: (a) GP-1B metallographic grinding and polishing machine; (b) Leica DM ILM metallographic microscope; (c) FALCON 450 Vickers hardness tester; (d) JB-W300A microcomputer-controlled pendulum impact testing machine.

2.2.2. Mechanical Characterization

Microhardness distribution was evaluated on both as-received and heat-treated specimens using Vickers microhardness testing. These tests were conducted in compliance with the ASTM E384 standard [32]. Measurements were taken across the clad steel, including the substrate, clad layer, and bi-material interface, using a FALCON 450 Vickers hardness tester (INNOVATEST Europe BV, Maastricht, The Netherlands) equipped with a diamond Vickers indenter. A load of 1 kg was applied with a dwell time of 10 s, as illustrated in Figure 2c.

Charpy V-notch impact testing was conducted to assess toughness evolution before and after heat treatments, following microhardness characterization. Standard specimens with dimension of 55 × 10 × 10 mm³ were prepared with a 2 mm deep 45° V-notch, machined on A240 TP316L-clad layer (Figure 3a), A238 Gr C parent metal (Figure 3b), and on the interface side (Figure 3c). Charpy impact tests were performed using a JB-W300A microcomputer-controlled pendulum impact testing machine (Jinan Testing Equipment IE Corporation, Jinan, China). Tests were performed at ambient temperature, as depicted in Figure 2d, with recorded absorbed energies.

Figure 3. Charpy standard specimens prepared with 45° V-notch machined on (a) the A240 TP 316L-clad layer; (b) the A283 Gr C carbon steel substrate, and (c) the interfacial side.

2.3. Experimental Design Using the Taguchi Method

The Taguchi method proved to be a reliable and efficient approach for the optimization of heat treatment parameters. This approach utilizes orthogonal arrays to reduce the number of experiments while maintaining the ability to analyze the effect of multiple variables. An L9 (3³) orthogonal array was selected, enabling the evaluation of three factors at three levels each using only nine experimental runs. The selected factors were:

-

A: Holding temperature (°C);

-

B: Holding time (min);

-

C: Heating and cooling velocity (°C/h).

The factors and their corresponding levels are presented in Table 2. The Taguchi design matrix is shown in Table 3.

Table 2. Heat treatment parameters and corresponding levels used in the Taguchi design.

Parameters	Level 1	Level 2	Level 3
Holding temperature (°C)	450	500	550
Holding time (min)	30	45	60
Heating and cooling velocity (°C/h)	100	150	200

Table 3. Taguchi L9 orthogonal array for the stress-relief annealing experiments.

Experiment	Holding Temperature (°C)	Holding Time (min)	Heating and Cooling Velocity (°C/h)
1	450	30	100
2	450	45	150
3	450	60	200
4	500	30	150
5	500	45	200
6	500	60	100
7	550	30	200
8	550	45	100
9	550	60	150

The mechanical property evaluated was toughness (Y), measured using Charpy impact tests. The response was modeled additively, considering only the main effects of each factor (Equation (1)), as follows:

$$\mathrm{Y}=\mathrm{\mu }+\mathrm{A}\mathrm{ᵢ}+\mathrm{B}\mathrm{ⱼ}+\mathrm{C}\mathrm{ₖ}+\mathrm{\epsilon }$$

(1)

where

-

$\mathrm{Y}:$ Measured toughness (in J/cm²);

-

$\mathrm{\mu }$: Overall mean of the response;

-

${\mathrm{A}}_{\mathrm{i}}$, ${\mathrm{B}}_{\mathrm{j}}$, ${\mathrm{C}}_{\mathrm{k}}$: effects of the selected level i, j, and k of each factor;

-

$\mathrm{\epsilon }$: Experimental error.

3. Results and Discussion

3.1. Characterization of the A283 Gr C/A240 TP 316L-Clad Steel Before Heat Treatment

3.1.1. Microstructural Characterization Before Heat Treatment

The microstructural characterization of the hot-rolled A283 Gr C/A240 TP 316L-clad steel was conducted to identify the structural features of the base metal (A283 Gr C), the cladding layer (A240 TP 316L), and the bi-material interface (Figure 4, Figure 5 and Figure 6). The optical micrographs of A283 Gr C carbon steel revealed a ferritic–pearlitic structure (Figure 4a), whereas the A240 TP 316L stainless steel exhibited a uniform, polygonal austenitic grain structure, as shown in Figure 4b.

Figure 4. Optical micrographs showing the as-received microstructures of the clad steel components: (a) ferritic–pearlitic microstructure of the A283 Gr C carbon steel; (b) fully austenitic microstructure of the A240 TP 316L stainless steel.

Figure 5. Optical micrographs showing the decarburized ferritic zone on the A283 Gr C carbon steel side near the interface: (a) general view; (b) higher-magnification view of the decarburized region.

Figure 6. Optical micrographs showing the carburized zone in the A240 TP 316L stainless steel near the interface: (a) general view; (b) higher-magnification view of the carburized region.

Microscopic examination of the bi-material interface revealed the presence of microstructural heterogeneities (Figure 5 and Figure 6). As shown in Figure 5b, the zone near the interface on the A283 Gr C steel side exhibits a coarse-grained ferritic structure with almost no pearlite, attributable to carbon depletion resulting from carbon diffusion from the substrate toward the clad layer. This decarburized ferrite zone has a measured thickness of 240.7 µm.

In contrast, on the A240 TP 316L stainless steel side, a carbon-enriched region, commonly referred to as a carburized zone, is observed near the interface (Figure 6b). These observations are in accord with the findings of [14,33,34]. It is worth noting that no diffusion of substitutional alloying elements such as chromium or nickel from the austenitic stainless steel into the carbon steel was detected (Figure 6). Consequently, no microstructural transformation or secondary phase formation was observed within the austenitic structure of the A240 TP 316L-clad layer. This indicates that during the hot-rolling process, the diffusion of alloying elements from the stainless steel toward the carbon steel substrate is negligible. This reflects the high structural stability of the austenitic phase as reported by Di Schino et al. [8].

3.1.2. Mechanical Properties

Microhardness characterization across the bi-material interface was conducted to evaluate the mechanical gradient between the A240 TP 316L stainless steel-clad layer and the A283 Gr C carbon steel substrate. The x-axis denotes distance from the interface (0 μm), with negative values in the stainless steel cladding and positive in the carbon steel substrate (Figure 7a).

Figure 7. Microhardness characterization across the bi-material interface: (a) microhardness profile showing the transition from A240 TP 316L stainless steel to A283 Gr C carbon steel; (b) optical micrograph illustrating the spacing between indentation points.

The non-heat-treated hardness profile reveals significantly higher hardness values in the cladding layer (200–240 HV) compared to the carbon steel substrate (105–130 HV). A sharp hardness gradient is observed at the bi-materaial interface, reflecting the contrast between the austenitic microstructure of the stainless steel cladding and the ferritic–pearlitic structure of the base metal. In the bi-material interface, a relatively noticebale increase is noted on the austenitic side within the carburized region, whereas a pronounced decrease in hardness is observed on the carbon steel side, corresponding to the decarburized ferritic zone. This behavior is primarily attributed to carbon diffusion during the hot-rolling, leading to carbon depletion in the ferritic zone near the interface.

To ensure the reliability of the measurements, appropriate spacing between indentation points was maintained to prevent interaction effects and localized strain hardening, as illustrated in Figure 7b.

To establish a baseline understanding of the as-received material, Charpy V-notch impact tests were conducted on non-heat-treated specimens in accordance with ASTM E23 standards [35]. The notches were machined at three distinct locations: along the interface side, within the A240 TP 316L stainless steel-clad layer, and in the A283 Gr C carbon steel substrate (as illustrated in Figure 3). For each notch position, three replicate tests were performed at room temperature, and the average impact toughness averages are reported in Table 4. To facilitate consistent comparative analysis across subsequent heat-treated conditions in this study, the notch positioned at the carbon steel–austenitic stainless steel interface was selected as the reference configuration. This approach captures the critical interfacial behavior that is most relevant to the overall mechanical integrity of the clad plate.

Table 4. Charpy impact toughness of the A283 Gr C/A240 TP 316L-clad steel before heat treatment with different notch locations.

	Toughness Average (J/cm2)
Notch on the interface side	161
Notch on A240 TP 316L steel	108.75
Notch on A283 Gr C steel	173.75

3.2. Characterization of the A283 Gr C/A240 TP 316L-Clad Steel After Heat Treatment Using the Taguchi Experimental Design

In this study, an L9 orthogonal array from the Taguchi method was employed, incorporating three key factors: holding temperature (450–550 °C), holding time (30–60 min), and heating/cooling velocity (100–200 °C/h). This design resulted in nine experimental conditions, as detailed in Table 3. The effects of these heat treatment parameters on metallurgical microstructure, Vickers microhardness profiles, and on the Charpy impact toughness of the A283 Gr C/A240 TP 316L-clad steel were systematically investigated. This is for the purpose of identifying optimal annealing conditions that maximize toughness, moderate interfacial hardness, and preserve structural integrity without inducing deleterious phases or excessive diffusion.

3.2.1. Microstructure Evolution

Following the stress-relief annealing defined by the Taguchi method (Table 3), detailed microscopic examinations were performed to evaluate the microstructural response of the clad material and to validate the selected processing parameters. The analysis focused on assessing microstructural modifications within the A240 TP 316L/A283 Gr C-clad steel system.

Figure 8 shows representative micrographs of the A283 Gr C carbon steel substrate after the different heat treatment conditions. The observations indicate that the applied thermal cycles do not induce significant microstructural changes in the carbon steel, as no secondary phase formation or abnormal grain growth was detected. Consistently, the microstructural analysis reveals a stable ferrite–pearlite structure throughout the investigated heat treatment conditions. These results confirm that the selected heat treatment parameters are effective in preserving microstructural stability.

Figure 8. Optical micrographs showing the microstructure of A283 Gr C carbon steel after stress-relief annealing processes, illustrating the ferritic and pearlitic phases.

Microscopic examination of the A240 TP 316L austenitic stainless steel-clad layer revealed no detectable microstructural changes or precipitation of secondary phases after stress-relief annealing (Figure 9). These observations are consistent with the microstructure prior to heat treatment (Figure 4b) and indicate that the applied thermal cycles did not induce phase transformations or undesirable diffusion phenomena. This is in agreement with the findings of Chu et al. [15], who explained that precipitation of intermetallic phases such as the chromium- and molybdenum-rich sigma (σ) phase in 316L stainless steel generally begins to appear at temperatures above 600–650 °C and depends strongly on exposure time and alloy composition.

Figure 9. Optical micrographs showing the microstructure of the A240 TP 316L stainless steel after different stress-relief annealing processes, confirming the stability of the austenitic structure.

In the present study, although annealing temperatures approached 550 °C, over relatively short holding times, the stability of the austenitic A240 TP 316L-clad layer effectively limited transformations [7,8,10,15]. This is supported by the absence of σ-phase or other precipitates in the microstructural observations. These results confirm that the selected stress-relief annealing conditions (450–550 °C, 30–60 min, 100–200 °C/h) preserve the microstructural stability and corrosion resistance of the A240 TP 316L-clad layer while avoiding the formation of deleterious phases.

3.2.2. Microhardness Measurement

Figure 10 displays the Vickers microhardness (HV) profiles across the A283 Gr C/A240 TP 316L-clad steel interface for the nine Taguchi L9 conditions, compared to the non-heat-treated baseline.

Figure 10. Microhardness profiles across the bi-material interface for different heat treatment conditions: (a) Experiment No. 1, (b) Experiment No. 2, (c) Experiment No. 3, (d) Experiment No. 4, (e) Experiment No. 5, (f) Experiment No. 6, (g) Experiment No. 7, (h) Experiment No. 8 and (i) Experiment No. 9.

After stress-relief annealing, the A240 TP316L stainless steel remained stable, with no evidence of carbide precipitation or σ-phase formation (Figure 6). In addition, stability of the microhardness values was noticed (Figure 10). These observations indicate the absence of secondary hardening and confirm that no embrittlement occurred during the thermal cycles [15].

The carbon steel A283 Gr C exhibits marginal softening (5–15% microhardness reduction), most evident within 180–640 μm of the interface line and pronounced in Experiments 7, 8, and 9 conducted at a holding temperature of 550 °C (Figure 10g–i). This decrease indicates that stress-relief annealing at this temperature successfully mitigated residual stresses via dislocation reduction and partial recrystallization [36].

These findings suggest the selection of 550 °C as the first parameter for the optimization of heat treatment. The holding temperature of 550 °C has demonstrated its effectiveness in refining the mechanical performance and stress distribution of the clad steel system.

3.2.3. Charpy Impact Toughness Measurement

The response variable in the Taguchi method was toughness, measured for each combination of holding temperature, holding time, and heating/cooling velocity, as summarized in Table 5. This experimental design allowed for the evaluation of the influence of heat treatment parameters on toughness while minimizing the number of experimental runs required. The specimens were prepared with 45° V-notch, machined on the interface side (Figure 3c).

Table 5. Taguchi L9 orthogonal array showing heat treatment conditions, measured Charpy impact toughness, and corresponding signal-to-noise (S/N) ratios.

Experiment	Holding Temperature (°C)	Holding Time (min)	Heating and Cooling Velocity (°C/h)	Toughness (J/cm2)	S/N Ratio (dB)
1	450	30	100	157.50	43.95
2	450	45	150	163.75	44.29
3	450	60	200	165.00	44.35
4	500	30	150	167.5	44.47
5	500	45	200	152.50	43.67
6	500	60	100	155.00	43.81
7	550	30	200	165.00	44.35
8	550	45	100	177.50	45.00
9	550	60	150	165.00	44.35

In accordance with the Taguchi methodology, the signal-to-noise (S/N) ratio was evaluated using the “larger-is-better” criterion (Table 5). This approach was selected because the objective of the study was to maximize the Charpy impact toughness. The S/N ratio was calculated using the following expression (Equation (2)):

$${\displaystyle \frac{\mathrm{S}}{\mathrm{N}}}=-10{\mathrm{l}\mathrm{o}\mathrm{g}}_{10}({\displaystyle \frac{1}{\mathrm{N}}}{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\displaystyle \frac{1}{{{\mathrm{y}}_{\mathrm{i}}}^2}})$$

(2)

where

${\mathrm{y}}_{\mathrm{i}}$: Measured toughness value;

$\mathrm{n}$: Number of experiences.

As shown in Table 5, the toughness values range from 152.5 to 177.5 J/cm², indicating a strong sensitivity of impact performance to the selected heat treatment parameters.

Among the nine experimental conditions, Experiment 8 exhibits the highest Charpy impact toughness (177.5 J/cm²), accompanied by the maximum S/N ratio (45.00 dB). This result demonstrates that the combination of a holding temperature of 550 °C, a holding time of 45 min, and a heating/cooling velocity of 100 °C/h provides the most favorable conditions for enhancing impact toughness. Under these conditions (Experiment 8), stress-relief annealing is particularly effective in reducing residual stresses and improving the material’s energy absorption capacity. In contrast, the lowest toughness value (152.5 J/cm²) is observed in Experiment 5, suggesting that intermediate temperatures (500 °C) combined with higher heating/cooling velocity (200 °C/h) may be less effective in enhancing impact toughness.

The simultaneous increase in toughness and S/N ratio with holding temperature indicates that temperature is the dominant parameter governing stress relaxation and microstructural recovery at the interface. The agreement between high toughness values and elevated S/N ratios demonstrates the robustness and reliability of the optimal heat treatment condition identified using the Taguchi method.

3.3. Statistical Analysis and Optimization of Heat Treatment Parameters

For the optimization of stress-relief annealing conditions, a detailed statistical analysis of the experimental data using main effects plots, interaction plots, and contour plots was conducted. These plots are completed to visualize and quantify the variations in Charpy impact toughness as a function of heat treatment conditions (Figure 11, Figure 12 and Figure 13).

Figure 11. Main effects plots showing the variation in average Charpy impact toughness as a function of (a) holding temperature, (b) holding time, and (c) heating and cooling velocity, with the highest toughness indicated by a yellow circle.

Figure 12. Interaction plots showing the combined effect of (a) holding temperature and holding time and (b) holding temperature and heating/cooling velocity, with the yellow circle marking the parameter combination yielding the highest toughness.

Figure 13. Contour plots showing the variation in Charpy impact toughness as a function of (a) holding temperature and holding time and (b) heating/cooling velocity and holding time. Color gradients represent constant toughness levels.

Figure 11 presents the main effects plots of the average Charpy impact toughness as a function of holding temperature, holding time, and heating and cooling velocity. The dashed horizontal line represents the average toughness of the non-heat-treated specimen notched on the interface side (161 J/cm²), which serves as a reference baseline. As shown in Figure 11a, the holding temperature has a pronounced effect on impact toughness. Increasing the temperature from 450 to 500 °C results in a slight decrease in toughness (158.3 J/cm²). Whereas a further increase to 550 °C leads to a marked improvement, yielding the highest average toughness value (169.2 J/cm²). This trend highlights the dominant role of holding temperature in promoting effective stress relaxation and microstructural recovery at the interface, thereby enhancing energy absorption under impact loading [36]. Figure 11b illustrates the influence of holding time. Charpy impact toughness increases with holding time up to 45 min. Beyond this, a decline is observed at 60 min. This behavior suggests that an intermediate holding time is optimal. Insufficient time limits stress relaxation, while excessive exposure may reduce the beneficial effects due to over-recovery or the redistribution of residual stresses [36].

The effect of heating and cooling velocity is shown in Figure 11c. A moderate velocity of 150 °C/h yields the highest toughness, whereas faster velocity (200 °C/h) results in lower values. This indicates that moderate thermal cycles (100 and 150 °C/h) may hinder uniform stress relief, while excessively rapid thermal cycles velocity (200 °C/h) provide no additional mechanical benefit.

The interaction plots further reveal that the beneficial effects of holding time and heating/cooling velocity are strongly dependent on the holding temperature (Figure 12). In particular, a holding temperature of 550 C consistently produces superior toughness, regardless of the secondary parameter considered. This confirms that holding temperature governs the response, while holding time and heating and cooling velocity act as secondary parameters.

Contour plots (Figure 13) provide additional insight into the combined influence of the heat treatment parameters on impact toughness. Regions of high toughness (>170 J/cm²) are predominantly concentrated around a holding temperature close to 550 °C (Figure 13a), a holding time of approximately 45 min (Figure 13a,b), and heating/cooling velocity in the range of 100–110 °C/h (Figure 13b). These contour maps clearly delineate a narrow and well-defined processing window in which impact toughness is maximized.

Beyond identifying a single optimal condition (Figure 11), the contour plots emphasize the existence of a robust parameter domain in which high toughness levels are maintained despite moderate variations in the processing conditions (Figure 13). This characteristic highlights the effectiveness of contour analysis for practical process optimization as it enhances the reliability of the heat treatment by reducing sensitivity to minor parameter fluctuations. Consequently, the contour plots provide strong support for the optimization of heat treatment parameters.

In addition to the statistical evaluation based on interaction plots and contour plots, an analysis of variance (ANOVA) was conducted to quantitatively assess the contribution of each heat treatment parameter. The outcomes of this analysis are presented in Table 6.

Table 6. ANOVA results used to optimize stress-relief annealing parameters based on Charpy impact toughness.

Source of Variation	DoF	Sum of Squares (SS)	Mean Square (MS)	Contribution (%)
Holding temperature	2	2.84	1.42	48.3
Holding time	2	0.29	0.15	4.9
Heating/cooling velocity	2	0.74	0.37	12.6
Total	6	—	—

The ANOVA results clearly indicate that holding temperature is the most influential factor, accounting for 48.3% contribution of the total variation in impact toughness. This dominant contribution confirms that temperature plays a critical role in governing stress relaxation and interfacial microstructural stabilization during stress-relief annealing [29]. The high mean square (MS = 1.42) associated with holding temperature further supports its strong statistical significance. In contrast, holding time exhibits a much lower contribution of only 4.9%, suggesting that variations within the investigated range (30–60 min) have a limited effect on toughness compared to temperature. Extending the holding time does not substantially improve energy absorption. The heating/cooling velocity contributes 12.6% of the total variation, indicating a non-negligible influence on impact toughness. This effect can be attributed to the residual stress development during heating and cooling. Excessively rapid velocities may reintroduce stresses, while moderate velocities favor mechanical stability [28].

Overall, the ANOVA results are in strong agreement with the trends observed in the interaction plots (Figure 12), contour plots (Figure 13), and signal-to-noise ratios of the Taguchi approach (Table 5). They confirm that holding temperature is the dominant parameter, followed by heating/cooling velocity, while holding time has a relatively minor effect.

These statistical analyses validate the selection of 550 °C as the optimal holding temperature, combined with an intermediate holding time (45 min) and minor heating/cooling velocity (100 °C/h), for maximizing interfacial toughness without inducing detrimental microstructural changes. These parameters correspond to the conditions applied in Experiment 8.

To validate the optimization of heat treatment parameters, a detailed microstructural examination and quantitative measurement of the decarburized ferrite zone were conducted on the specimen from Experiment 8. The results show a clear reduction in the thickness of the decarburized ferrite zone (Figure 14), with a decrease of approximately 80 µm compared to the non-heat-treated clad steel condition (Figure 5b). This substantial reduction decreases the thickness of mechanically weak regions at the interface, thereby mitigating interfacial embrittlement and enhancing the overall mechanical integrity of the clad steel. This is in agreement with the findings reported by Liu et al. (2018) [14].

Figure 14. Optical micrograph showing the reduction in the thickness of the decarburized ferrite zone in Experiment 8 after optimized stress-relief annealing.

In conclusion, the mechanical behavior of the hot-rolled A283 Gr C/A240 TP 316L-clad steel is strongly influenced by the dislocation structure generated during the rolling and cladding process. In the as-rolled condition, a high dislocation density results in elevated internal stresses and localized strain accumulation. This condition leads to increased hardness but reduces impact toughness due to stress concentration near the interface [7,14,16,36].

The stress-relief annealing applied in Experiment 8 promotes recovery and partial recrystallization within the carbon steel substrate. During this process, dislocations rearrange into lower-energy configurations, reducing the overall dislocation density. As a result, internal stresses are relaxed, and microhardness decreases in the decarburized ferrite zone (Figure 10h). This stress relaxation mitigates strain localization at the interface and enhances the material’s capacity to absorb impact energy. The improvement is reflected by the increase in Charpy impact toughness (Table 5). In addition, the reduction in the thickness of the decarburized ferrite zone limits the extent of mechanically weak regions at the interface (Figure 14). Consequently, interfacial embrittlement is reduced, contributing to improved overall mechanical integrity of the clad steel.

Therefore, the results clearly demonstrate the effectiveness of this stress-relief annealing treatment in improving the mechanical integrity of the A283 Gr C/A240 TP 316L-clad steel, preserving microstructural integrity (Figure 8 and Figure 9).

4. Conclusions

This study successfully optimized the stress-relief annealing parameters for hot-rolled A283 Gr C/A240 TP 316L stainless steel-clad plates using the Taguchi design methodology. The combined experimental and statistical approach enabled a clear identification of the heat treatment conditions that maximize impact toughness while preserving microstructural stability.

• Microstructural analyses confirmed that stress-relief annealing within the investigated temperature range does not alter the ferritic–pearlitic structure of the carbon steel substrate, nor the austenitic structure of the 316L-clad layer, and no evidence of carbide precipitation or secondary phase formation was detected. A reduction in the thickness of the decarburized ferrite zone at the interface was observed, indicating improved interfacial integrity under optimized conditions.
• Mechanical characterization revealed a moderate decrease in microhardness near the interface, accompanied by a significant improvement in Charpy impact toughness. These changes were attributed to recovery and partial recrystallization phenomena in the carbon steel substrate, leading to reduced dislocation density, relaxation of residual stresses, and enhanced energy absorption during impact loading.
• Statistical analysis demonstrated that holding temperature is the most influential parameter governing toughness, followed by heating/cooling velocity, while holding time has a comparatively minor effect.
• The optimal stress-relief annealing conditions were identified as a holding temperature of 550 °C, a holding time of 45 min, and a heating/cooling velocity of 100 °C/h.

Overall, this work demonstrates that stress-relief annealing, when properly optimized, is an effective strategy for improving the mechanical performance and interfacial reliability of hot-rolled stainless steel-clad plates, providing valuable guidance for industrial applications.