*5.3. Ferritic Ductile Irons*

Among all the ductile iron structures, LST in ferritic DIs has gained importance in the last decades because it lacks a significant pearlite amount (whose carbon content is higher due to the presence of cementite plates). This means that it is very difficult to find conventional ways of transforming the matrix to a martensitic type, and, on the other hand, it is difficult to achieve hardening without melting, since carbon diffusion from graphite nodules locally lowers the melting point, favoring the occurrence of melting near the surface [39].

The first modeling of the phase transformations during LST of ferritic DIs can be found in Grum and Sturm's study [63], where LSM was carried out with a low-power CO2 laser. To ensure the fusion of the surfaces, an overlapping of 30% was considered. Three new regions arise after LSM: the melted zone, the hardened zone, and an intermediate transition zone. In the region where the liquidus temperature is exceeded, the predominantly ferritic matrix melts, and the graphite nodules diffuse toward the melted surface. This causes a strong dissolution of carbon in the liquid matrix, whose rapid solidification transforms it into austenite and ledeburite dendrites. The intermediate zone is characterized by a highly localized fusion process around the graphite nodules that, depending on the magnitude of the carbon diffusion toward their edges, can generate a ledeburite and/or martensite ring. The transformation scheme is shown in Figure 13. The differences between the transition zone and the hardened zone are that, although the matrix transforms into austenite during the heating cycle and is enriched in carbon due to diffusion away from the graphite nodules, this is not enough to locally reduce the melting point, and, therefore, the matrix becomes of martensitic type with residual austenite. In addition, due to the heterogeneous distribution of the nodules, there are areas with low carbon content where austenite reverts to ferrite during cooling.

**Figure 13.** Schematic representation of phase changes around graphite nodules during LSM: (**a**) Initial DI microstructure, (**b**) transformation to nonhomogeneous austenite during heating, (**c**) carbon enrichment in austenitic matrix after diffusion, (**d**) local melting around graphite nodules, (**e**) formation of martensitic shells around graphite nodules after cooling, (**f**) formation of ledeburitic shells above graphite nodules that suffered local melting (Reprinted with permission from ref. [63]. 1996 Elsevier).

In Grum and Sturm [64], the analysis of LSM in ferritic DIs is completed, experimentally and numerically comparing the thickness of the martensite and ledeburite layers that surround the graphite nodules in transition and hardening zones. As described in Roy and Manna [57], from simplified models of temperature during the heating and cooling cycle, in conjunction with diffusion equations based on Fick's law, the authors determined, at different depths, if the temperature around the nodule is enough to produce fusion (and the formation of ledeburite when cooling), in addition to calculating the magnitude of the radial diffusion of the carbon atoms. As shown in Figure 14, the simulation precisely adjusted to what was experimentally determined, which validates the use of the aforementioned models. The average standard deviation obtained is less than 10% and can be explained by model simplifications, mainly in terms of diffusion, as average values that do not vary with temperature were used.

**Figure 14.** Comparison of the experimental and predicted thickness of martensite and ledeburite shells around graphite nodules (Reprinted with permission from ref. [64]. 2002 Elsevier).

This microstructural analysis was completed in Grum and Sturm [65], where the authors applied LSM to a similar DI sample to verify their previous models, as well as XRD and residual stress tests. The metallic surface was characterized by a high content of martensite (41%) and a portion of residual austenite (23%), and at higher depths, a reduction of the martensitic phase was observed in favor of the austenitic phase (24% and 42%, respectively, at 240 μm from the surface). From these results and the residual stresses measured at different points on the surface, it was determined that the residual austenite and martensite contents were the most influential on the nature of the internal stresses (Figure 15). The stresses on the surface always have a tensile character, since the transformations to austenitic phases involve a volume reduction of such magnitude that they cancel the effect of the transformation to martensite, thus implying a volumetric increase and, consequently, stresses of a compressive type.

**Figure 15.** Evolution of residual stresses at the ductile iron surface as a function of volumetric % of phase constituents (reprinted from Reference (Reprinted with permission from ref. [65]. 2005 Inderscience Enterprises Ltd.).

Benyounis et al. [66] presented a new approach on the surface melting of DIs by conducting a comparison between the microstructural changes caused by two technologies: Nd:YAG laser and TIG. In the case of LSM, the laser was used at low power (0.1 kW) and speed of 1 mm/s with 50% overlapping, while in the case of TIG, the voltage was maintained at 50 V and the current was varied between 80 to 120 A to ensure fusion. The microstructure obtained by LSM consists of a fine dendritic structure of retained austenite, with residual percentages of martensite and cementite (Fe3C). Although in TIG melting a dendritic structure was also observed in an austenite and cementite matrix, the main difference was the higher amount of retained austenite in LSM, which was attributed to a superior cooling rate.

The dendritic microstructure is also present in Alabeedi et al. [67], where LSM was carried out on an 83% ferrite, 11% pearlite, and 6% graphite ductile iron. A Gaussian distribution CO2 laser was used at 3 kW power, with a scanning velocity of 10 mm/s. To achieve melting, a 50% overlapping was considered, and argon shielding gas was applied to also prevent oxidation and contamination of the specimen. The obtained microstructure was identified by SEM as austenite dendrites immersed in a cementite network, attributed to the high cooling rate during solidification. An additional transition region was observed below the melted zone; it was formed by a mixture of dendrites and thick martensite plates and added to some undissolved nodules surrounded by ledeburite layers.

Recent studies have further supported these results; for instance, Pagano et al. [68] carried out a complete analysis of the effects of the LSM technique on a ferritic DI, using an

Nd:YAG laser operating at 1 kW of power, with circular spot and Gaussian distribution. A parabolic transformed cross-section was obtained, coinciding with the laser distribution. In the area where fusion occurred, the microstructure became a network of austenite and ledeburite dendrites, with a small portion of martensite. Then a transition region with localized fusion was observed around the graphite nodules that did not melt. Finally, a hardened area was appreciated, where the graphite spheres are surrounded by a layer of martensite and immersed in a matrix of ferrite and residual austenite. These results corroborate the work of Grum and Sturm [63,64].

Another approach for LSH/LSM was conducted by Catalán et al. [69], where a fiber delivery diode laser with uniform energy distribution and rectangular shape was used. Two linear power ramps, along with two scanning velocities, were defined to achieve four different sets of linear energy densities in order to establish a relationship between the experimental results and the joint effect of all laser parameters, as well as to determine the transition between hardening from solid-state and melting. It was concluded that, for highvelocity settings (1000 mm/min), microstructural changes were not significant, whereas, for low-velocity cases (570 mm/min), severe changes could be observed. The melted zone depth increased directly proportional with the average linear energy (Figure 16), correlating with changes in the linear energy absorbed by the material and the thermal diffusion properties of the surface-modified samples. The chemical composition, as determined by GDOES, was strongly affected by the laser, with a significant increase for C and decrease for Si compared to the as-cast sample, due to the transformation of graphite nodules to carbides during the heating process. Regarding input parameters, at low energies, a Fe3C phase was detected by XRD, whereas γ-Fe2O3 was raised at high energy densities, suggesting that surface oxidation occurred during the heating cycle.

**Figure 16.** (**a**) Melted zone (MZ) and hardened zone (HZ) of laser surface-modified microstructure of DI; (**b**) melted depth as a function of the average linear energy of the laser (Reprinted with permission from ref. [69]. 2021 Springer Nature BV).

On the subject of hardness improvement, many of the reviewed works have quantitatively addressed the efficiency of LSM in DIs. For instance, Grum and Sturm [63] observed a uniform hardness distribution throughout the transformed layer (800–900 HV0.1), and when the limit of the hardened zone ( ≈550 μm) was exceeded, it fell rapidly to the base value (around 250 HV0.1), and this is in agreemen<sup>t</sup> with the posterior measurements of Pagano et al. [68]. Benyounis et al. obtained a harder microstructure when using TIG melting over LSM (750 and 500–600 HV, respectively), since the supplied energy and the interaction time ensured less retained austenite content, implying that the tempering effect of melting was reduced.

In the melted zone, Catalán et al. [69] found a stable microhardness of 1000–1100 HV0.3, which was five times higher than the nominal value. This was attributed to the carbide/oxide formation in the transformed matrix, as observed in Figure 17, since the mentioned XRD analysis supported the existence of these harder phases. The main difference

with other works is the uniformity in both hardness measurements and microstructure for the melted and hardened zones, since the laser energy distribution was uniform and rectangular shaped. Thus, overlapping of laser tracks was not needed and the tempering effect was avoided.

**Figure 17.** (**a**) Evolution of microhardness as a function of depth and linear energy. (**b**) Surface hardness and composition of as-cast and laser modified DIs (Reprinted with permission from ref. [69]. 2021 Springer Nature B.V.).

The wear response of laser melted DIs is broadly detailed in Pagano et al.'s [68] study, where a rotating steel cylinder for sliding wear tests was chosen to roll against the samples before and after being treated with LSM. The evolution of the coefficient of friction and material loss was recorded as a function of the total sliding distance. In this test, two stages were observed: initially, the coefficient of friction (COF) increases linearly with the sliding distance (run-in stage), and then it decreases to a stationary value (steady-state), due to the lubricating effect of graphite release. In this phase, since there is a partial dissolution of graphite in the hardened area, an increase in the adhesive component of the COF is generated by reducing the lubricating effect. Furthermore, the presence of hard phases, such as cementite and martensite, causes an increase in the abrasive component of the COF, which accounts for doubling the friction coefficient of the untreated DI. From Archard's equation for sliding wear, we obtain the following:

$$W = K \left(\frac{F\_N \cdot S}{H}\right)'\tag{9}$$

the authors verified that, for the same sliding distance (*S*), applied load (*FN*), and wear factor (*K*), the greater hardness (*H*) of the LSM sample generated a lower volumetric loss *V*. This was supported by the fact that the original samples showed rapid plastic deformation during the test, while the presence of oxides on the surface of the specimens modified by LSM gave evidence of a moderate triboxidative wear regime.

Subsequently, Ceschini et al. [70] compared the tribological experiments carried out by Pagano et al. [61] for four different specimens: untreated GI (GJL300), as-cast DI (GJS400), low-energy melted DI (GJS400-LHV at 500 W), and DI treated with LSM at high energy (GJS400-HHV at 1000 W), using the same equipment and tests as the work in comparison. Due to different energy densities, the two LSM-treated DIs exhibited different surface hardness, with the GJS400-HHV being the hardest (1100 HV1, compared to 850 HV1 for the GJS400-LHV). The sample with the best performance was the GJS400-LHV, followed by the GJS400-HHV, as observed in Figure 18. In both cases, a slight triboxidative wear regime was observed, such that the differences reside in the presence of microcracks in the HHV cast iron, associated with its greater hardness and lower toughness. Furthermore, the results showed that graphite morphology in cast irons is a major factor in wear resistance, with the lamellar graphite shape (GJL300) being the one that acts as the greatest stress concentrator, favoring nucleation and crack growth during the sliding test.

**Figure 18.** Worn surfaces of different cast irons for a 20 N applied load (Reprinted with permission from ref. [70]. 2016 Elsevier).

The erosion resistance enhancement due to LSM has been also assessed by Alabeedi et al. [67] in their study, where silica particles were bombarded at an average speed of 50 m/s at different incident angles. The amount of mass loss as a function of exposure time increased linearly, as in Gadag and Srinivasan [55], with a cumulative loss at the end of each trial (45 min) between 35 and 100 times less in the case of the treated samples. Furthermore, SEM examination evidenced that the wear mechanism in the untreated specimens begins with the deformation of the softer graphite nodules, which then favor the nucleation of cracks by acting as stress concentrators. Thus, the cracks form larger craters as they spread, so that the surface fractures and becomes dimpled. In contrast, wear in laser-modified samples is attributed to the initiation of fatigue cracks, which develop only into small craters. However, the hard structure formed in the process does not allow these cracks to propagate through the matrix and cause more damage, mainly due to the presence of retained austenite.

Finally, Boccardo et al. [71] developed a unidirectional coupled thermo-metallurgical model in order to predict the phase transformations of laser-treated DIs. In particular, the thermal model computed the temperature evolution during the process, while the metallurgical model predicted the phase evolution as a function of the temperature. The possible transformations during heating include the reverse eutectoid transformation (RET), homogenization of the austenite carbon content (HA), and melting transformation (MT), whereas the cooling process can comprise the eutectoid transformation (ET), solidification (ST), and martensitic transformations (MDT and MWT). To test the performance of this model, the authors considered the same DI as in Catalán et al. [69], with the four sets of linear power ramps and scanning velocities detailed in their work. Figure 19 shows the relationship between experimental and simulation results for 180 J/mm energy density. According to the presented metallurgical model, the observed thickness of the melted zone (Figure 19a) agrees with the computed temperature (Figure 19c), because it is greater than the temperature at which melting transformation ends (*TLe*). For a depth of 100 μm, the material partially melted, because the temperature remains between the limits where melting transformation starts (*TLs*) and ends. Finally, for depths bigger than 300 μm, there is no phase transformation, since the temperature does not reach the point where RET starts. On the other hand, Figure 19b illustrates the evolution of the volumetric fraction of ledeburite, one of the main constituents of ductile iron during these transformations. The comparison between the computed ledeburite fraction and the thickness of melted, transition (hardened), and base material zones (Figure 19a) shows that the thermo-metallurgical model is able to identify these regions as a function of the ledeburite fraction in the transformed matrix. A linear correlation between measured and simulated martensite/ledeburite thickness layers was obtained, thus verifying that, for low thicknesses, the thermo-metallurgical model reasonably captures the trend of experimental results, and it slightly overestimates the thickness at higher magnitudes. Moreover, the thickness of the ledeburite–martensite

layer depends on the laser parameters, and it is increased with the increment of laser power and the decrement of scanning velocity.

**Figure 19.** (**a**) Optical microscopy image of resulting microstructure of modified DI. (**b**) Computed ledeburite volume fraction at fixed laser parameters. (**c**) Computed temperature distribution at a fixed instant and parameters as a function of depth (Reprinted with permission from ref. [71]. 2021 Springer Nature B.V.).

### *5.4. Overall Facts of Surface Laser Treatment of Cast Irons*

Considering the literature review, some similarities can be found in LST of cast irons, regardless of their initial microstructure. As presented in Figure 20, laser modification of iron castings generally leads to a martensitic or ledeburitic microstructure, depending on the amount of supplied energy and the subsequent phase transformations.

**Figure 20.** Summarizing diagram of microstructural transformations and overall response of lasertreated cast irons.

Moreover, mechanical/tribological characterization shows an improvement on hardness and wear resistance after LST, as explained by the formation of harder phases during rapid solidification. Microhardness curves tend to exhibit a stable, harder region, with a quasi-exponential decay to the as-cast value in zones with low degree transformation. In the case of tribological tests, mass loss is significantly reduced by both LSM and LSH; however, the wear rate is highly variable, and it depends on the conditions of the friction system. Finally, since LST encompasses volumetric transformations, the analysis of residual

stresses is important, and it usually reveals a change from compressive stresses near the surface to tensile stresses further into the as-cast region.
