Microstructural characterization of AISI 431 martensitic stainless steel laser-deposited coatings

Figure 1 shows the three distinct microstructures which could be identified in the deposits. Observation of interfacial regions (Fig. 1a) showed that a thin layer of planar growth was present at all interfaces with an extension of around 20 μm. This plane growth zone became unstable as solidification proceeded and broke down into cellular dendrites within short distances from the interface. As a result, cellular dendrites formed the majority of solidification microstructures. In multi layer deposits, a columnar to equiaxed transition (CET) occurred at the top of deposits resulting in a thin layer of equiaxed grains (Fig. 1b).

To explain the observed microstructural changes, it is necessary to know the parameters which control the solidification structure. These parameters are mainly solid–liquid interface growth rate R and temperature gradient in the melt G. The G/R ratio determines the solidification mode and the G · R product controls scale of the solidified structure [9]. CET takes place when nucleation of equiaxed grains occurs in the liquid ahead of the columnar structure. In laser cladding, this will not happen unless the ratio of G/R is lower than some critical value [10]. As can be seen in Fig. 2, by lowering the G/R ratio, solidification structure changes from planar to cellular and then columnar and equiaxed dendrites. In single layer coatings, only the transition from planar to cellular and columnar dendrites happened but in five layer deposits, all of the transitions were visible. These changes imply that during solidification, G/R ratio was not constant. Previous research [4] shows that in laser solidification treatments the G/R is infinite at the bottom of the melt pool which translates into a planar growth layer but rapidly decreases and as a result, structure of laser claddings mostly consists of columnar growth of dendrites [4].

Recent studies [11, 12] confirmed that in multi layer laser claddings with constant power, temperature, and size of the melt pool increase from each layer to the next because of heat accumulation. This was also observed in this study in which size of the melt pool in multi layer deposits (measured by optical microscopy after sectioning) was almost doubled from the first to the fifth layer. In fact, this heat accumulation reduces the high temperature gradients typically associated with columnar microstructures and provides the opportunity for a transition from columnar to equiaxed structure [11].

Increasing the cladding speed (i.e., increasing the solidification rate) did not affect the microstructural features such as epitaxial growth across subsequent layers or transitions in the solidification mode but scale of the dendritic structures was significantly refined as is obvious by the reduction of cell spacing at higher cladding speeds as shown in Fig. 3. Higher cladding speeds produced smaller cell spacing with values down to around 9 μm in five-layer and 3 μm in single layer deposits. This can be explained by the fact that growth rate is a direct function of cladding speed and higher growth rates are obtained at higher deposition velocities. A variation of 15% is included in the data of Fig. 3 to account for changes in cell spacing values from bottom to top of the coating because of local variations in solidification conditions [13]. The difference in the cell spacing values of single and multi-layer deposits in spite of their similar cladding speeds can be attributed to the heat accumulation in multi-layer claddings as explained above but in both cases, cell spacing followed the empirical power law of cell spacing V ^−0.5 reported previously [14].

The connection between chemical composition of stainless steel laser depositions and their equilibrium phase constitution can be obtained by equilibrium diagrams such as Schaeffler and ternary phase diagrams. Figure 4 shows the composition of AISI 431 on Schaeffler diagram [15] and on a section of ternary Fe–Cr–Ni phase diagram at the latest stage of solidification (1,400 °C) calculated by 3D Demo system for Ternary Phase Diagrams from Japan Science and Technology Corporation. A dilution of around 10% from substrate was taken into account in calculating the position of AISI 431 on the above-mentioned diagrams but because of big differences in the amounts of C, Cr and Ni between substrate and coating, Cr_eq. and Ni_eq. values with and without dilution were approximately the same and this allowed neglecting dilution from substrate and having the same position on diagrams for both single and multi layer deposits.

Based on these diagrams, equilibrium solidification of AISI 431 is a mixed austenitic-ferritic solidification which finally results in around 10% delta ferrite (δ) along with martensite and retained austenite at room temperature. While these diagrams are reliable for predicting structures at cooling rates close to equilibrium, numerous previous researches [3, 16‐19] showed that solidification and solid state cooling rates have a significant effect on the final microstructures, thus making it unreliable and even impossible to predict the phase constitution of rapidly cooled stainless steels from the conventional diagrams. This unreliability was shown in the current research in terms of the amount and distribution of δ in claddings deposited at different speeds. Figure 5 shows the optical micrographs of 5-layer samples deposited with cladding speeds of 5 (Fig. 5a) and 117 mm/s (Fig. 5b). It can be seen, especially at track overlapping areas, that by increasing the cladding speed less δ (visible as white islands) was formed. In addition, distribution of δ is not homogenous and as shown in Fig. 6, while little amounts of δ are present at interdendritic areas; huge quantities are formed at the beginning of each solidification step, i.e., at substrate-coating interface or in remelted parts of the previous track. A point which should be emphasized here is that high cooling rates after solidification tends to suppress the solid state transformation of δ to γ. Consequently, concentration of δ in some areas should be a result of solidification step.

Studying the solute redistribution in the overlapping areas (which contain considerable quantities of both γ and δ) by EDS spectral mapping showed that even in the lowest cladding speed, there was no detectable partitioning of alloying elements between these two phases. Figure 7 shows the EDS maps for Cr, Ni (obtained at 20 kV), and C (obtained at 4 kV) as the main alloying elements in AISI 431 stainless steel. Uniformity of chemical composition in dual phase areas suggests that solidification of these coatings occurred in a partitionless manner probably because of the high cooling rate and its associated undercooling during the laser deposition process.

Taking the high cooling rate and partitionless solidification into account, the liquidus line in phase diagram will not be the start of solidification anymore and instead, T _o line should be considered as the beginning of solidification as shown in Fig. 8. T _o line is the loci of T _o temperatures at which the new phase can appear with a net decrease in free energy at the same composition as the parent phase. In fact, T _o line defines the boundary below which diffusionless transformation is possible. Below T _o, free energy of the single phase γ or δ is lower than the free energy of undercooled liquid and partitionless solidification will be theoretically possible [3]. T _o curve is located between the liquidus and solidus lines of equilibrium phase diagram (Fig. 8) in which the data from references [3] and [18] are combined to show the composition of AISI 431 in the phase diagram with T _o lines. In this case, solidification structure depends on whether the T _o line (or its extension) of γ, δ or both of them is crossed by the undercooled melt and relative quantity of γ and δ is governed by the difference in T _o temperatures of the two phases [16]. As can be seen in Fig. 8, T _o temperatures of γ and δ for AISI 431 composition are very close and this means that if the undercooled liquid crosses the T _o lines, both phases will form in similar quantities. This idea is confirmed by micrographs of Fig. 6 which show that at the beginning of each solidification stage, γ and δ phases were formed in similar proportions. It can be concluded that the undercooling achieved during rapid solidification was sufficient to bring the liquid metal below both T _o lines making it possible for both phase to nucleate.

Figure 6 also shows that after a short time, γ outgrew δ and became the primary solidification phase. This can be attributed to the faster growth kinetics of γ in comparison to δ which was observed in several previous research works on welding of stainless steels [3, 16, 20]. It seems that while partitionless solidification provides the opportunity for both γ and δ to nucleate at initial stages of solidification, more favorable growth kinetics of γ (especially at higher solidification rates) dominates the rest of solidification resulting in a mostly austenitic structure with isolated block of δ at the interfaces and occasional stringers at interdendritic areas.

It is expected that by changing the cladding speed, not only the solidification rate changes, but also cooling rate in solid state is affected and this can have implications on solid state transformation products. According to TTT diagram of the AISI 431 stainless steel [21], this type of steel has a very high hardenability and to form any transformation product other than martensite, such slow cooling rates are needed that take 10³–10⁴ s to reach around 600 °C from austenitizing temperature. Under our processing conditions, having these slow cooling rates is virtually impossible and only the transformation of austenite to martensite can occur. Optical microscopy and SEM microstructural investigation of the samples deposited at different cladding speeds revealed that this was the general case in which typical lath martensite was the dominant microstructural feature. Figure 9 shows examples of martensitic microstructure of the coatings.

Orientation imaging microscopy (OIM) was employed to analyze the characteristics of solid state transformation products with special interest in austenite/martensite orientation relationship (OR) and characteristics of the austenite/martensite interphase boundaries. Figure 10 shows examples of grain map, phase map and [001] inverse pole figure (IPF) map of the studied coatings which are superimposed on image quality (IQ) map (The IQ parameter describes the quality of an electron backscatter diffraction pattern [22]). In these figures, martensite lathes and traces of retained austenite are clearly visible. The data of the images presented in Fig. 10 are used in the subsequent calculations of OR and boundary characteristics.

As the formation of martensite involves coordinated movement of atoms, parent austenite and product martensite always have some form of orientation relationship (OR). Over the years, several models such as Kurdjumov–Sachs (K–S) [23], Nishiyamam–Wassermann (N–W) [24, 25], and Greninger–Troiano (G–T) [26] have been proposed to describe the fcc/bcc OR formed during martensitic transformation. The basic idea in all of these models is that closed-packed planes of the two phases (i.e., {111} for austenite and {110} for ferrite) are parallel or nearly parallel. By confirming the validity of the concept of parallel close-packed planes, it will be reliable to use these models to study the OR in our coatings. To explore this task, pole figure (PF) texture plots for close-packed planes of austenite and martensite were obtained and as shown in Fig. 11, it is obvious that the criterion of parallel close-packed planes is valid in our case because the projection of all {111} plane poles for austenite coincides with the projection of some {110} planes for martensite.

Having this fundamental concept confirmed, discrete PF plots for different planes of austenite and martensite were obtained and their combined PF plot was compared with the theoretical PFs for each OR model. The combined PF plot of the two phases shows close resemblance to the PF of G–T model as shown in Fig. 12.

Usually different austenite–martensite ORs are associated with different values of misorientation across the boundaries between these two phases [27]. Consequently the misorientation across the austenite–martensite boundaries was also studies by OIM to check the existence of G–T orientation relationship. Figure 13 shows a close-up of an area of the images in Fig. 10 containing both martensite and retained austenite.

This section was selected in such a way that it contained only one originally austenite grain (evident from the inverse pole figure map in Fig. 13b) to avoid considering the γ–γ interfaces in the boundary measurements. By measuring the misorientation angle across interphase boundary, it was known that the two phases are separated by high angle boundaries with misorientation of around 45°. More detailed analysis of the misorientation values showed that there was in fact a range of boundary misorientations from 42 to 46 degrees as highlighted in Fig. 13d. Highlighting the 42–46° boundaries in the whole scan area revealed that this is the general case and almost all of the interphase boundaries are high-angle boundaries with misorientations in the range of 42–46° (Fig. 14). Average misorientation angle between austenite and martensite was 44.27° which is very close to 44.3° misorientation calculated for G–T orientation relationship from Eulerian angles typical of this OR [27]. By repeating the above-mentioned analyses for single and multi layer coatings deposited at the lowest (5 mm/s) and the highest (117 mm/s) cladding speeds, it became obvious that different cooling rate associated with various cladding speed had no noticeable effect on boundary characteristics or OR in the final structure of the coatings.

On the other hand, EBSD results showed that while in multi layer deposits the amount of retained austenite was around 4% for all cladding speeds, in single layer coatings more austenite (up to 11%) was retained at the higher cladding speed. This results in a variation of hardness values of single layer deposits produced at different scanning speeds as shown in Fig. 15a but in multi layer coatings, hardness values were the same in all cases (Fig. 15b). The observed differences can be attributed to the austenite stabilizing effect of finer dendrites produced at higher cladding speeds [13].