*5.2. Austempered Ductile Irons*

Since austempering produces a unique initial matrix, LST can lead to different microstructural transformations and mechanical properties. Lu and Zhang [56] first carried out LSH with a CO2 laser on bainitic ADI samples. The material followed a standard austempering process, with an austenitizing temperature of 890 ◦C, a cooling interval at 360 ◦C for 2 h, and a final quenching to ambient conditions. The initial austenitic bainite matrix also exhibited ferrite and austenite enriched in carbon, with a smaller portion of retained austenite and martensite. After laser hardening, the matrix changed to an acicular structure of martensite and retained austenite, with graphite spots close to the surface, as seen in other ductile iron samples.

An extensive numerical model of LST in ADIs was carried out by Roy and Manna [57], with the goal to predict the temperature in the vicinity of graphite nodules. The estimation is based on the heat equation for the case of a continuous wave CO2 laser with Gaussian distribution profile. This equation is expressed as follows:

$$
\nabla^2 T - \frac{1}{\alpha} \frac{\partial T}{\partial t} + \frac{q\_r}{\lambda} = 0,
\tag{7}
$$

where *T* represents the temperature, α is the thermal diffusivity, *qr* is the magnitude of the heat delivered by the laser per unit volume and time, *λ* is the thermal conductivity, and *t* is the time. Assuming that heat losses are negligible and that the thermal properties of the material do not depend on the temperature, the equation was solved analytically, and the temperature profile was used to determine the magnitude of carbon diffusion away from the graphite nodules, according to Fick's law, written as follows:

$$\mathcal{C}(y,t) = \frac{\mathcal{C}\_f - \mathcal{C}\_t}{2} \left[ 1 - \text{erf}\left(\frac{y}{2\sqrt{Dt}}\right) \right] + \mathcal{C}\_{t\prime} \tag{8}$$

which determines the carbon concentration toward the edges of the nodule. In this expression, *Cf* is defined as the carbon% in the austenitic phase at the matrix–nodule interface, *Ce* represents the carbon concentration in the matrix, *y* is the measured horizontal distance from the center of the nodule (assuming a perfect circular shape), *D* is the diffusion coefficient (which depends on temperature), and *t* is the time. This carbon concentration was then utilized to define, from the iron–carbon–silicon (Fe–C–Si) phase diagram, what the temperature is when fusion occurs. Then, by selecting the proper laser parameters, this model would allow us to achieve the material hardening without melting. The authors conclude that, for laser powers below 800 W and a fixed scanning velocity of 60 mm/s, the fusion width is negligible compared to the distribution of nodules in the matrix, and, therefore, the microstructure is predominantly martensitic.

Subsequently, Roy and Manna [58] compared the results of LSH and LSM on ADI samples, using sets of parameters to ensure melting and hardening without melting. A variable laser power between 0.8 and 1.5 kW and a speed between 20 and 1000 mm/s were selected for the case of LSM, and a power between 0.5 and 1 kW and a speed of 60 mm/s were chosen for LSH analysis. The depth of the transformed zone maintained a linear relationship with the power and an inverse correlation with the scanning velocity, regardless of the chosen parameters. Microscopic observation and X-ray diffraction analysis (XRD) made it possible to determine that LSM produced a higher proportion of retained austenite in the metallic matrix, while, in LSH, martensite was the predominant phase.

Zammit et al. [59] also compared the results of implementing LSH and LSM on ADI, considering a discrete spot laser in order to avoid the reported hardness reduction in overlapping zones. A 9 kW CO2 laser with Gaussian energy deposition profile was used to generate stationary pulses, decreasing the total energy input and, thus, producing lower distortion and processing costs. First, it was determined that melting caused severe distortion on the surface, given by an increment in surface roughness from 0.15 to 1.34 μm (in arithmetic value). Moreover, the molten region was depleted of graphite nodules, forming soft phases, such as austenite and residual ferrite, along with low carbon martensite. On the other hand, LSH only increased the average surface roughness from 0.15 to 0.43 μm, while the microstructure was mainly composed of martensite with unaltered graphite nodules. Some nodules were surrounded by a bullseye ledeburitic structure, implying that the carbon diffusion lowered the melting point locally around them, as in Roy and Manna [57]. Soriano et al. [60] were then able to verify the microstructural features of LSH in ADI by using an Nd:YAG laser coupled with a PID controller to keep the temperature in the center of the laser constant at 975 ◦C, so that the melting point was not exceeded during the treatment, as in Liu and Previtali [41]. A hardened zone composed of coarse martensite in the near-surface, as well as a finer acicular structure below it, was distinguished, associated with the rapid solidification without melting of the material (Figure 10a). Moreover, as evidenced in Figure 10b, the mixture of retained austenite and upper and lower bainite structures below the coarse martensitic array verified the occurrence of a softer thermal cycle in this region.

**Figure 10.** Micrographs of (**a**) top and (**b**) medium zones of the laser hardened ADI sample (Reprinted with permission from ref. [60]. 2011 Elsevier).

Regarding the hardening effect of the LSH, the profile as a function of depth, measured by Lu and Zhang [56], with a Knoop test with 0.2 kg load, followed the trend of the results shown in Gadag et al. [50], with a homogeneous hardness of approximately 600 KHN in the transformed zone (~250 μm) and a subsequent rapid decay to the base value of the ADI specimen (300 KHN). Putatunda et al. [61] provided more evidence of the effects of LSH on the mechanical properties of ADIs. Hardness, yield stress, and ultimate tensile strength tests were conducted after initial austempering conditions similar to those of Lu and Zhang [56]. Changes in these mechanical properties were analyzed for four cases: (a) untreated ADI; samples treated with the same laser parameters, (b) initially without any type of coating; and then with two types of coating, namely (c) graphite and (d) manganese. LSH significantly increased most of the properties (except the ultimate tensile strength), obtaining the best performance under the use of coatings, since a greater energy absorption was responsible for a cooling rate that ensured the formation of a more resistant martensitic structure.

In Soriano et al. [60], the improvement of hardness was further backed up by the analysis of residual stresses (Figure 11) from XRD patterns, using the sin2(ψ) method. The diffraction peak position of the (211) α-Fe phase was measured at nine different inclinations, ranging from −45◦ to 45◦. The hardness of the heat-affected zone was quantified in the range of 700–800 HV, which decreased as the fraction of retained austenite increased in favor of the martensitic phase. Moreover, in the hardened region, compressive residual stresses were measured and explained by the volumetric increase associated with the transformation from austenite to harder martensite. As the depth increased, the degree of martensitic transformation decreased until the point where the matrix corresponds to the base material (0.9–1.5 mm), and, thus, the stresses become tensile. This result was also addressed by Zammit et al. [59] in their study, where it was found that compressive stresses take place near the surface, due to the 4% volumetric increase of austenite to martensite transformation. At approximately 160 μm, the stresses became tensile, due to the presence of retained austenite in the ausferritic bulk, therefore obtaining an estimation of the HAZ depth of the laser treatment.

The tribological performance of laser-modified ADIs was first assessed in Lu and Zhang [56], where the characteristics of sliding wear were analyzed from a 500-min test against a steel disc without lubrication, at a relative speed of 1.2 m/s and a variable load between 2 and 14 kgf. The results showed a linear relationship between the mass loss due to friction and the time of the test, as in Gadag and Srinivasan [55], but it was also determined that the wear rate, in both cases, consisted of two stages: a mild one (for loads less than 10 kgf) and a severe one (for loads greater than 10 kgf). These stages were related to microstructural changes during the wear test, where the austenite phase decreased as the applied load increased. The mild phase of wear is characterized by an oxidative mechanism near the surface, while the severe phase is dominated by lamination and delamination processes, the effect of which was observed by SEM in the formation of microcracks in the surface.

Roy and Manna [58] established a relationship between the hardness profiles and the wear resistance of ADIs. The hardness distribution in LSM was parabolic, while, for LSH, it was approximately constant within the transformed zone. Thus, the presence of a more homogeneous hardened zone in LSH provided the best tribological behavior, as observed in Figure 12. For a 5 kg load, three stages in the evolution of wear were identified. Initially, a high wear rate was associated with maximum contact area. Then a stationary stage was established, due to the adherence of surface asperities originated in the first phase, to finally give way to an accelerated wear stage due, to the formation of furrows and debris separation. The optimal performance obtained by LSH is supported by a lower extent of micro-fractures that is attributed to the lower probability of causing microcracks during friction contact. This result is in agreemen<sup>t</sup> with the deformation restrictions at the austenite–martensite interface that are caused by changes in the residual stresses during the transformation from austenite to martensite, as detailed in References [59,60].

**Figure 12.** Evolution of displaced material as a function of sliding wear distance for as-received and laser-treated ADI samples (left); SEM micrographs of worn surfaces (right) for (**a**) LSM treated ADI, (**b**) magnified view of wear debris of (**a**), and (**c**) LSH-treated ADI (Reprinted with permission from ref. [58]. 2001 Elsevier).

The tribological characterization of LSH was completed in Zammit et al. [62]. Based on the results of Reference [59], the authors carried out discrete spot LSH treatments by using the same CO2 laser at 600 W and 300 ms of pulse duration and considering three different arrays: laser spots separated by one spot diameter, adjacent spots, and 50% overlapped spots. Scuffing (pin-on-disc) and rolling contact fatigue (RCF) tests were performed to address the wear resistance and mechanisms. In the first test, as-austempered and laser-treated ADI pins were considered, with a hardened and oil-lubricated AISI D2 steel counterpart disc rotating at 1450 rpm under 10 MPa constant pressure. LSH treatment with adjacent spot tracks exhibited the highest sliding cycles to failure, followed by separated spots, showing that the back-tempering effect on hardness, due to overlapping, reduced the ADI wear resistance by an order of 10–100 times. Moreover, adjacent spot LSH treatment induced a higher martensite volume fraction, which is harder than the initial microstructure and is able to delay crack initiation and propagation. On the other hand, RCF tests were performed only on as-austempered and adjacent spot laser-treated ADIs. The LSH increased the number of cycles until fatigue failure over 10<sup>6</sup> times. This increment was also attributed to a harder microstructure and compressive residual stresses near the surface. It was concluded from SEM and pit observations that the main wear mechanism is governed by plastic deformation and propagation of cracks around graphite nodules, and the overall wear resistance is comparable to that of carburized steels. As a conclusion, the authors sugges<sup>t</sup> that laser-treated ADIs could be suitable replacements for a variety of engineering components.
