Effect of Austempering Time and Temperature on the Mechanical and Microstructural Properties of a Niobium-Alloyed Austempered Ductile Iron

Abstract

This study evaluated the influence of niobium addition and austempering time and temperature on the microstructure and mechanical behavior of ductile iron. Three alloys were produced: unalloyed ductile iron (H1) and two Nb-alloyed ductile iron (H2, 0.11 wt.% Nb and H3, 0.32 wt.% Nb). After austenitizing at 900 °C for 60 min, samples were austempered at 250 °C and 300 °C for 15, 30, 60, and 90 min. The as-cast microstructure of H3 exhibited a higher pearlite fraction (73.31 vol%) and increased carbide content (2.48 vol%), accompanied by reduced nodularity and nodule count. X-ray diffraction analysis revealed that the highest fraction of carbon-rich retained austenite was obtained in H3 after 30 min at 300 °C, reaching 42.48%. Hardness decreased with increasing retained austenite, confirming the inverse relationship between this phase and matrix strengthening. Wear testing showed that H2 presented slightly lower volume loss due to carbide precipitation, with the lowest value recorded after 15 min at 300 °C (1.088 mm³). Tensile tests indicated that ultimate tensile strength and yield strength were superior at 250 °C, with H3 achieving the highest values at 90 min (1353 and 1090 MPa, respectively). Overall, niobium promoted carbide formation and austenite stabilization, modifying the balance between hardness, toughness, and wear resistance in austempered ductile iron.

1. Introduction

Austempered ductile iron (ADI) is an alloy with excellent mechanical properties [1]. It has gained attention due to factors such as low cost, fatigue resistance, wear resistance, and its high strength-to-weight ratio [2]. ADIs are widely used across diverse industrial sectors, including the automotive, mining, railway, and agricultural industries, in components such as crankshafts, camshafts, suspension parts, and wear-resistant machinery elements. Compared to steels and other types of cast iron, ADIs offer superior machinability, and their processing often requires lower energy consumption than that of conventional alloys, thereby enhancing their potential as a sustainable engineering material [3].

Austempering is an isothermal heat treatment applied to ferrous materials to enhance their strength and toughness [4,5].

This heat treatment involves two steps. First, the ductile iron is heated to the austenitizing temperature, typically between 850 and 950 °C to achieve a fully austenitic matrix. Then, it is rapidly cooled to avoid pearlite formation, maintaining a temperature range of 250–400 °C. To minimize oxidation and decarburization effects, a salt bath furnace is recommended for ADI processing [6].

The strength of ADI depends on the morphology of the ausferritic matrix, which is influenced by the austempering temperature and time [6,7,8,9]. During austempering, two transformation reactions take place, where unstable austenite (γ) transforms into acicular ferrite and carbon-rich austenite [9,10,11]. In the first stage, austenite is transformed into acicular ferrite (α) and carbon-rich austenite (γ_HC), as shown in Equation (1) [12,13].

$$\mathsf{\gamma }\to a+{\gamma }_{HC}$$

(1)

In the second stage of austempering, carbon-rich austenite transforms into bainitic ferrite (α_Bain) and carbides such as Fe₃C or ε-carbides. This stage is undesirable because the carbides negatively impact the mechanical properties of the material, making it more brittle. This stage is represented in Equation (2) [14,15].

$${\gamma }_{HC}\to {\alpha }_{Bain}+Carbides$$

(2)

The time interval between the end of the first stage and the beginning of the second stage is known as the optimal process window, where the best balance of mechanical properties is achieved [16,17].

Retained austenite with high carbon content plays a fundamental role in the mechanical behavior of ADIs. Its presence significantly enhances ductility and strength, while also improving toughness and fatigue resistance. The amount and stability of retained austenite are primarily controlled by austempering temperature and time. At low austempering temperatures (250–300 °C), a fine ausferrite microstructure develops, characterized by a higher volume of fine acicular ferrite and a low fraction of retained austenite. This is due to slower austempering kinetics and the distribution of carbon into austenite in very thin films between ferrite needles during the early stages of austempering. These thin films have limited capacity to form and stabilize large volumes of austenite, and atomic mobility is reduced at these temperatures, restricting carbon diffusion and retained austenite formation [18,19].

At higher austempering temperatures (310–400 °C), carbon atomic mobility increases, facilitating its diffusion into austenite. In addition, the coarser and more widely spaced acicular ferrite morphology provides greater space for retained austenite to form and stabilize in larger volumes [20].

Several alloying elements have been incorporated into ADIs to modify their microstructure and mechanical properties, including Mo, Al, Ni, and Cu. For molybdenum, Olanrewaju et al. [21] reported that increasing Mo content (0.11 to 0.24 wt.%) increased the pearlite fraction (49.38 to 56.59 vol%) and influenced damping behavior. Colín et al. [22] added 0.1 and 0.3 wt.% Mo and austempered at 270 °C for holding times of 60, 90, and 120 min, finding that higher Mo contents (0.3 wt.%) promoted carbide formation (3.91 vol%) and resulted in the highest fraction of high-carbon retained austenite (11.5 vol%).

Regarding aluminum, Adebayo et al. [23] studied different Al contents (1.05, 1.57, 2.29, 3.02, and 3.74 wt.%) at austempering temperatures of 300, 350, and 400 °C, reporting that Al addition induced ferrite precipitation around graphite nodules and that austempering improved both tensile properties and hardness. In the case of copper, some studies [24,25] have shown that Cu decreases strength but increases elongation and impact energy, as it promotes the formation of high-carbon retained austenite, particularly at higher temperatures. Finally, nickel has also been shown to affect mechanical properties; one study [26] reported that Ni increased impact resistance without significantly changing ultimate tensile strength, while another study [27] observed a gradual increase in tensile strength with Ni content up to 2 wt.%.

The use of niobium (Nb) in ADI is relatively recent compared to its applications in steels [28]. Fras et al. [29] studied the effect of small amounts of niobium (above 0.038 wt.% Nb) on the microstructure and mechanical properties, finding that niobium increased the nodule count while reducing nodule size. Souza et al. [30] reported that a 0.47 wt.% Nb addition resulted in a slight improvement in tensile ductility and Charpy impact energy. Similarly, Chen et al. [31] observed a significant effect on graphite morphology, noting that nodularity decreased from 92.4 wt.% to 84.5 wt.% when the Nb content increased from 0 to 0.11 wt.%. Chen et al. [32] also evaluated the effect of niobium in ADI. They found that NbC nanoparticles act as nucleation sites for graphite and ferrite needles, thereby increasing the volume fraction of ausferrite in the microstructural matrix of ADI. Finally, Abdullah et al. [33] reported that the mechanical properties of ADI with 0.25 wt.% Nb depend on austempering time. Despite these studies, comprehensive knowledge of ADI with niobium remains limited.

This work investigated the effect of niobium microalloying additions on the fraction of high-carbon retained austenite in ADI at low austempering temperatures (250 and 300 °C) and different austempering times (15, 30, 60, and 90 min). Three material systems were studied: H1, corresponding to ADI without niobium addition, H2, corresponding to ADI with 0.11 wt.% niobium and, H3, corresponding to ADI with 0.32 wt.% niobium. These systems were selected to evaluate how the incorporation of niobium influences the microstructure and mechanical properties of ADI. The main objective is to quantify the retained austenite fraction using X-ray diffraction and to analyze its effect on the hardness, toughness, and wear resistance of the three systems.

2. Materials and Methods

The production of unalloyed ductile iron (H1) and niobium-alloyed ductile irons (H2 and H3) was carried out using an induction furnace without protective atmosphere, starting with ASTM 1018 steel (DISA, Hidalgo, Mexico), ferrosilicon (25.4 mm × 9.6 mm, 75 wt.% Si), and graphite (4.76 mm × 3.18 mm, 99 wt.% C). Niobium was added in the form of ferroniobium (25.4 mm × 9.6 mm, 65 wt.% Nb). Subsequently, inoculation was performed with FeSi, and nodularization was carried out using FeSiMg (12.7 mm × 9.6 mm, 7 wt.% Mg) by pouring the molten material at 1450 °C into a cast iron mold and allowing it to cool to room temperature.

The chemical composition of the ductile irons was determined using a Shimadzu spark spectrometer PDA-7000 (Shimadzu Corporation, Kyoto, Japan), averaging five measurements for each alloy. The chemical composition of each casting is presented in

Table 1. Chemical composition of ductile irons (% mass fraction).

Element wt.%	H1	H2	H3
C	3.76	3.62	3.67
Si	2.54	2.64	2.61
Mn	0.38	0.32	0.32
S	0.011	0.010	0.009
P	0.009	0.007	0.009
Mg	0.047	0.046	0.047
Nb	0	0.11	0.32
Cr	0.02	0.02	0.01
Ti	0.002	0.004	0.003
Mo	0.01	0.01	0.01
Ni	0.046	0.042	0.044
Cu	0.22	0.19	0.23
CE	4.60	4.50	4.54

. For the calculation of the carbon equivalent (CE), Equation (3) [34] was employed, which requires the weight percentages of carbon (X_C), silicon (X_Si), and phosphorus (X_P).

$$CE=X_C+{\displaystyle \frac{1}{3}}(X_{Si}+X_P)$$

(3)

The correct amount of silicon can increase nodularity and enhance the mechanical properties of thick sections while inhibiting coarse graphite formation [35,36]. Low manganese content is typical in ductile irons due to its strong segregation in intercellular regions, where it promotes the formation of iron or alloyed carbides [37]. To ensure proper graphite nodule formation in ductile iron, the residual magnesium content must remain between 0.02 wt.% and 0.05 wt.% [38]. The volume and size of carbides increase significantly in the presence of niobium, which also alters the morphology of nodular graphite, reducing nodule count and nodularity [39]. The material was then cast into ingots with the dimensions shown in Figure 1.

The heat treatment consisted of two stages. In the first stage, the samples were austenitized in a muffle furnace at 900 °C for 60 min to ensure a fully austenitic phase in all specimens. Subsequently, they were austempered in a salt bath composed of 50 wt.% KNO₃ and 50 wt.% NaNO₃ at temperatures of 250 and 300 °C, to evaluate the effect of niobium at low austempering temperatures and its impact on the mechanical properties. Austempering was performed with holding times of 15, 30, 60, and 90 min, thereby determining the completion of the first austempering reaction and the onset of the second stage. Finally, the samples were quenched in water at room temperature. The process is schematically illustrated in Figure 2.

For metallographic analysis, transverse cuts were made in the central part of the ingot, marked as segments AB and CD, as shown in Figure 1. Specimens measuring 15 mm × 10 mm × 10 mm were extracted, ground with silicon carbide papers of grit sizes 50, 100, 200, 400, 600, 1000, 1500, and 2000, and subsequently polished with 1-micron diamond paste. Metallographic etching was performed using 3 vol% nital to reveal ausferrite and pearlite, and 10 vol% ammonium persulfate to reveal carbides. Microstructural characterization was carried out using two optical microscopes: a Keyence VHX-X1 microscope was employed for high-magnification imaging, and a Nikon 550 microscope was utilized for low-magnification observations. Three specimens were prepared for each material tested. Nodularity was calculated using Equations (4) and (5) [40], where Equation (4) provides the degree of graphite sphericity (SSF) based on the area (A) and perimeter (P) of nodules with diameters greater than 10 μm.

$$SSF={\displaystyle \frac{{\displaystyle 4·\pi ·A}}{P^2}}$$

(4)

Nodules with a sphericity greater than 0.65 were considered accepted nodules and their area (A_A) was measured. Conversely, nodules with a sphericity lower than 0.65 were classified as unacceptable nodules, and their area (A_Un) was obtained. Both values were used in Equation (5) to calculate the nodularity percentage (%Nod).

$$\%Nod=\left({\displaystyle \frac{A_A}{A_A+A_{Un}}}\right)·100$$

(5)

For the nodule spacing (λ_G), which indicates the carbon diffusion distance Equation (6) was employed [40], where NS_avg corresponds to the average nodule size, calculated using Equation (7), in which A represents the nodule area with diameters greater than 10 μm.

$${\lambda }_G=55.4{\left({\displaystyle \frac{{NS}_{avg}}{P^2}}\right)}^{{\displaystyle \frac{1}{3}}}$$

(6)

$${NS}_{avg}=2\sqrt{A{\pi }^{-1}}$$

(7)

Nodule counting was performed using an image analyzer, quantifying nodules larger than 10 µm in diameter and excluding those located at the micrograph boundaries. A total of eight different micrographs were analyzed. Finally, to quantify the percentages of carbides, pearlite, and ferrite, a threshold filter was applied in the image analyzer.

X-Ray diffraction (XRD) analysis was performed using an Inel Equinox 2000 diffractometer (INEL, Artenay, France) with Co-kα₁ radiation to determine the volume fraction of carbon-rich austenite (V_y). The procedure followed the methodologies of Putatunda et al. and Bedolla et al. [41,42], requiring data on intensities, diffraction angles, and wavelengths, as applied in Equation (8):

$$Vy={\displaystyle \frac{1-Vc}{1+\left(\frac{\mathsf{\Sigma }I\alpha }{\mathsf{\Sigma }Iy}\right)\left(\frac{\mathsf{\Sigma }Ry}{\mathsf{\Sigma }R\alpha }\right)}}$$

(8)

where Vc is the percentage of other phases (carbides and graphite), Iα and Iγ are the intensities of the hkl plane of the α and γ phases, respectively, and Rα and Rγ are values calculated according to Equation (9), where v is the volume of the unit cell, F is the structure factor times its complex conjugate, p is the multiplicity factor of the (hkl) reflection (i.e., the number of crystallographic planes that produce reflections of the same intensity due to symmetry), θ is the Bragg angle and e^−2M is the Debye–Waller or temperature factor, where M is a function calculated using Equation (10), with B being a material dependent constant related to the Debye temperature of the solid and λ the wavelength of the radiation source.

$$R=\left({\displaystyle \frac{1}{v^2}}\right)\mathrm{ }\left[{\left|F\right|}^{2}p(1+{\mathrm{c}\mathrm{o}\mathrm{s}}^{2}2\theta )/\left({\mathrm{s}\mathrm{i}\mathrm{n}}^2\theta \mathrm{c}\mathrm{o}\mathrm{s}\theta \right)\right]\left(e^{-2M}\right)$$

(9)

$$M=B{\displaystyle \frac{{Sin}^2\theta }{{\lambda }^2}}$$

(10)

Mechanical testing included impact, hardness, wear, and tensile evaluations. The Charpy impact test was performed on specimens in accordance with ASTM E23 [43], using a 45° V-notch with a depth of 2 mm. Results were obtained from three repetitions at room temperature.

Hardness was measured in accordance with ASTM E18 [44] using a Buehler Rockwell hardness tester (Model LAM/ME-2; Buehler, Lake Bluff, IL, USA) on the Rockwell C scale, employing a diamond indenter and a 1471 N loading force; ten indentations were performed on polished sample faces to determine average values. Wear resistance was assessed using a TE 53 SLIM Multi-Purpose Friction and Wear Tester (Phoenix Tribology Ltd., Newbury, United Kingdom) in accordance with ASTM G132-96 [45], under a load of 45 N and a speed of 300 rpm for 4000 revolutions, with three measurements per specimen averaged for reporting. Finally, tensile tests were performed on an Instron Series 3400 universal testing machine following ASTM E8/E8M [46], and the results of elongation, ultimate tensile strength (UTS), and yield strength (YS) correspond to the average of four tests per alloy.

3. Results and Discussion

3.1. Metallographic Analysis of Ductile Iron As-Cast Conditions

In Figure 3, the unetched microstructures reveal the presence of graphite nodules. It is essential to analyze graphite morphology, including parameters such as nodularity, nodule count, inter nodule distance, and nodule size.

Table 2. Constituent distribution in ductile irons.

Propety/Alloy	H1	H2	H3
Nodule size (µm)	21.72 ± 1.65	29.96 ± 1.32	28.73 ± 1.27
Nodularity (nodularity%)	87.23 ± 1.57	81.50 ± 1.73	81.90 ± 1.69
Internodule distance (µm)	23.92 ± 1.19	35.51 ± 1.62	30.66 ± 1.32
Nodule count (nod/mm2)	256.1 ± 27.56	113.7 ± 23.18	156.4 ± 25.12
Carbides (vol%)	0.53 ± 0.07	0.95 ± 0.05	2.48 ± 0.07
Pearlite (vol%)	25.44 ± 0.31	63.17 ± 0.22	73.31 ± 0.28
Ferrite (vol%)	66.03 ± 0.17	23.42 ± 0.23	17.54 ± 0.19

summarizes the analysis of these nodular features. A decrease in nodularity is observed, dropping from 87.23 nodularity% to 81.9 nodularity%. This reduction is attributed to the tendency of niobium to form carbides and to increase cementite stability during solidification, which suppresses graphitization and promotes the formation of irregular graphite [47]. Similarly, nodule count also decreases in the niobium-added alloys. Nodule count is closely related to the number of nucleation sites available in the melt; the addition of niobium promotes the formation of NbC carbides that act as competitive phases during nucleation [48], thereby reducing the number of active sites for graphite nodules. As a result, the inter0nodule distance increases compared to the unalloyed sample. Figure 3 also shows graphite nodules within a ferrite–pearlite matrix, as observed in micrographs etched with 3 vol% nital. Table 2 further indicates an increase in pearlite content due to niobium addition. This behavior can be explained by the segregation of niobium into interdendritic regions, where it precipitates as NbC. This precipitation consumes carbon, which decreases graphitization and instead favors the growth of carbon-enriched austenite that, upon cooling, transforms into pearlite (ferrite + cementite) rather than into graphite and ferrite [49], thereby producing a higher volume fraction of pearlite and a lower fraction of ferrite. For carbide analysis, 10 vol% ammonium persulfate etching was used. The white regions in correspond to carbides, as confirmed by the quantitative results in Table 2.

3.2. Microstructural Analysis of Austempered Ductile Iron

Figure 4, Figure 5 and Figure 6 display the metallographies of the three alloys after austempering heat treatment, revealing an ausferritic matrix. Figure 4 shows the ADI H1 samples at different temperatures and times, while Figure 5 and Figure 6 show the ADI H2 and ADI H3 samples, respectively, also under different temperatures and times, in the three cases the dark regions correspond to acicular ferrite, while the white regions indicate carbon-rich austenite. A high nodule count improves the ausferritic microstructure and shortens the optimal process window [50].

3.3. X-Ray Diffraction Analysis

Analyzing the diffractograms in Figure 7, the H3 alloy exhibits a higher percentage of carbon-rich austenite, as niobium acts as a stabilizing element for this phase. This effect is caused by the increased volume fraction of NbC precipitates in the interdendritic regions, as shown in Table 2, which serve as barriers to carbon diffusion, leading to carbon accumulation in the austenitic matrix. As a result, a greater fraction of carbon-enriched austenite remains stable at room temperature, increasing the amount of retained austenite. At 300 °C, the percentage of carbon-rich austenite further increases, as shown in Figure 8, due to the formation of coarse ausferrite, which contains a higher proportion of retained austenite enriched in carbon. This behavior has also been reported in previous studies [2]. Moreover, a clear relationship is observed between austenite content and material hardness: higher percentages of retained austenite result in lower hardness values due to the increased ductility associated with the austenitic phase.

3.4. Impact Testing

Figure 9 presents the results of the Charpy impact test. In the as-cast condition (0 min), H1 exhibits higher impact energy compared to H2 and H3. This reduction is associated with the higher proportion of hard and brittle phases such as pearlite and NbC carbides derived from the niobium addition, which act as preferential sites for crack nucleation and propagation. In addition, the decrease in nodularity and the increase in interparticle spacing reduce the effectiveness of graphite nodules in deflecting crack paths, facilitating a more continuous and rapid crack propagation [51]. After austempering heat treatment, impact energy values improved, with the highest value observed at 90 min and 300 °C in the H2 (13.8 J). This enhancement is attributed to improved energy dissipation in the nodule-ausferrite phase [52] and to the presence of a higher volume fraction of carbon-rich austenite, which increases impact resistance, as reported in other studies [14].

All the irons show lower impact energy at 250 °C, attributed to the formation of fine ausferrite at this temperature. Fine ausferrite results in a higher volume fraction of acicular ferrite, which is more brittle and dissipates energy less effectively.

3.5. Hardness Testing

The hardness test results shown in Figure 10 indicate that, in the as-cast condition, H3 exhibits higher hardness than H1 and H2. This is attributed to the greater amounts of carbides and pearlite in H2. It has been reported that niobium addition [27,29] increases carbide formation, thereby enhancing hardness due to two main factors: first, carbides act as reinforcing particles that hinder dislocation movement, and second, niobium reduces the interlamellar spacing of pearlite, generating more interfaces that act as obstacles to dislocation motion. The highest hardness after austempering was observed in H1 austempered at 250 °C for 15 min (40.76 HRC), as lower austempering temperatures lead to a higher volume fraction of acicular ferrite, which increases hardness [53,54]. Conversely, Figure 10 also shows that hardness is lower in austempered H2 samples. This behavior is attributed to niobium accelerating ausferrite formation during the austempering stage [55].

3.6. Wear Testing

The specimens austempered at 300 °C exhibited greater wear compared to those treated at 250 °C, as shown in Figure 11. This behavior can be attributed to the higher fraction of high-carbon retained austenite, which, due to its lower hardness, is more susceptible to plastic deformation and the formation of microcracks under sliding contact conditions. At the microscopic level, wear is governed by the interaction between harder acicular ferrite plates and softer austenitic regions: while ferrite contributes strength and toughness, an excessive amount of retained austenite promotes localized deformation and the initiation of subsurface cracks, accelerating material loss. Similar findings have been reported in previous studies [22,56], where higher austempering temperatures produced coarser ausferrite and higher fractions of austenite, both factors contributing to reduced wear resistance. Furthermore, the untreated H1 specimen exhibited a wear volume of 5.032 mm³ compared to 2.85 and 0.79 mm³ for H2 and H3, respectively. This difference is associated with a higher ferrite fraction and a lower carbide content in H1, which reduces its resistance to abrasive and adhesive wear mechanisms.

3.7. Tensile Testing

The results of ultimate tensile strength (UTS) are presented in Figure 12. It can be observed that UTS shows a marked increase compared to ductile iron in the as-cast condition. Ductile irons austempered at 250 °C exhibit higher UTS values than those austempered at 300 °C. This behavior is attributed to the formation of a fine ausferritic microstructure at lower austempering temperatures, which is characterized by a high-volume fraction of acicular ferrite and a lower volume fraction of carbon-enriched austenite. The combination of these microstructural features acts as an effective barrier to dislocation motion, thereby increasing the strength of the material [57,58].

Figure 13 shows the yield strength (YS) results. Higher austempering temperatures (300 °C) result in lower YS due to increased carbon diffusion, which stabilizes a larger fraction of high-carbon austenite and promotes the growth of acicular ferrite into coarser, less defective morphologies. As a result, the yield strength decreases compared to samples that austempered at 250 °C. The presence of niobium slightly modifies this behavior, as NbC carbides precipitate in the solid state, limiting the growth of acicular ferrite. Therefore, at lower austempering temperatures, the strengthening effect from grain refinement combines with the fine acicular morphology, leading to slightly higher YS values [59].

Elongation decreased significantly during austempering, as shown in Figure 14. In the as-cast condition, the highest values were observed in H1 (8.92 elongation%), which is attributed to the presence of ferrite and pearlite with a higher ferrite fraction, a soft and ductile phase. After 15 min of heat treatment, elongation decreased markedly, particularly at 250 °C, where the microstructure consists of fine ausferrite with a high defect density. This combination increases strength but substantially reduces ductility, Additionally, the presence of niobium slightly reduced elongation under the evaluated conditions. This behavior is associated with the precipitation of NbC carbides, which harden the matrix and limit the material’s ductility [60].

4. Conclusions

The addition of 0.11 and 0.32 wt.% Nb to ductile iron strongly influenced the microstructure and mechanical response after austempering at 250 and 300 °C. In the as-cast state, the Nb-alloyed alloy (H3) showed a markedly higher pearlite fraction (73.31 vol%) and carbide content compared to the unalloyed alloy (H1), which was associated with reduced nodularity and nodule count but larger nodule size. During austempering, niobium promoted the stabilization of carbon-rich austenite, reaching a maximum of 42.48 vol% at 30 min and 300 °C in H3. The retained austenite fraction was inversely correlated with hardness, whereas wear resistance improved in the Nb-alloyed austempered ductile irons due to the higher fraction of carbides. Tensile testing revealed that the finest ausferritic microstructures at 250 °C provided the highest strength, with H3 achieving an ultimate tensile strength of 1353 MPa and a yield strength of 1090 MPa after 90 min. Overall, niobium addition favors carbide formation and austenite stabilization, enhancing the strength–toughness synergy and offering a viable strategy to tailor austempered ductile iron for demanding structural and wear-resistant applications.