Influence of Process Conditions on the Local Solidification and Microstructure During Laser Metal Deposition of an Intermetallic TiAl Alloy (GE4822)

Figure 3 shows the temperature curve of a fixed spatial point during the build-up of a layer for an explanatory construction height of 5 mm.

The selected measurement point lies approximately 5 mm from the edge in the direction of application of the first of the 12 tracks that form a layer. The thermal influence of the adjacent tracks decreases with increasing distance from the measuring point (1st track). The applied material reaches temperatures of > 900 °C several times, which might significantly influence the microstructure formation in the solid phase, as can be seen from the CCT diagram for GE4822 elaborated in Reference 22. The authors found the amount of transformed α-phase to be highly dependent on the cooling rates by which the transformation range can be extended to temperatures even well below 900 °C for cooling rates > 10 °C/s, which applies to LMD. It is also noteworthy that the inert gas flow causes an additional cooling effect, which occasionally during the process cycle leads to temperatures below the preheating temperature.

Figure 4 shows the temperatures at the layer building start for various build heights for different beam diameters with a preheating temperature set to 900 °C.

With a beam diameter of 1 and 2 mm at a build height of 5 mm, maximum values of the measured starting temperatures are reached, and the temperatures drop when reaching 10 mm build height. With a beam diameter of 0.6 mm, on the other hand, the starting temperature drops continuously. The temperature increase for 1 and 2 mm beam diameters is due to the greater amount of heat introduced by higher laser power compared to the parameter set with 0.6 mm beam diameter. Additionally, the total process time for the large beam diameters is reduced by a factor of about 3 (d_s = 1 mm) respectively 5 (d_s = 2 mm). Since the inert gas settings for all processes are the same, the inert gas cooling is more important for slower processes, thus decreasing the starting temperature for d_s = 0.6 mm. Due to the smaller number of tracks (\( d_{s} \propto \) track width → allowing increased heights), the integrated melt pool surface through which heat can be dissipated fast is also smaller. Compared to this, the cooling effects (surface, shielding gas flow) predominate over the heat input in a setup for a beam diameter of 0.6 mm. This results in a decreasing preheating temperature with increasing build height.

The reduced preheating temperature in dependence on the build height is maximal at 10 mm. A possible explanation is the overall larger sample volume and small thermal diffusivity of TiAl compared to other metals (approximately 8.5 mm²/s). Additionally, temperature gradients perpendicular to the surface of the volume as well as between the surface and volume are likely to occur; in the context of induction, comparable effects are known as the sink effect.[23] The measured values displayed would thus represent a spatially distorted image of the temperature distribution in the volume.

In Figure 5, the temperature distributions from the FEM simulations for single melt tracks (virtually cut at symmetry) during LMD on similar preheated (to 900 °C) substrate material made of GE4822 are depicted.

Visible are solidified (blue-purple, T ≤ 1510 °C) and molten areas of the traces (red-yellow-white, T > 1510 °C). Both the different melt pool dimensions and exceeding the aluminum boiling temperature of 2450 °C for 1 and 2 mm beam diameter can be recognized. According to the calculations, the maximum melt pool temperature of 2905 °C is reached at a beam diameter of 1 mm. The reason for this is the intensity distribution, normalized to the maximum value shown in Figure 6. The average intensities are I = 93.9 W/mm² (d_S = 2 mm) I = 255 W/mm² (d_S = 1 mm) and I = 233 W/mm² (d_S = 0.6 mm). However, the graph (Figure 6) shows that the intensity in the center of the laser spot is much greater with a beam diameter of 1 mm than with diameters of 2 or 0.6 mm. Therefore, the melt pool temperature in the center of the spot is also higher.

For comparison of the solidification conditions, the solidification rate, temperature change rate or cooling rate and local temperature gradients are important. The third can be calculated from any two of the values. Figure 7 shows the simulated cooling rates at the melt pool boundary as a function of the position on the solidification front.

Cooling rates < 0 K/s depict heating of the material. The smaller the beam diameter, the higher the maximum cooling rate is. For a beam diameter of 0.6 mm, the cooling rate is in the range of 10⁴ K/s, whereas with a beam diameter of 2 mm, maximum cooling rates of the order of 10³ K/s are achieved. The cooling rates determined by analysis of the thermographic data obtained in this study range from 10³ to 10² K/s on the surface of solidified tracks and thus are consistent with the model. In Reference 21 cooling rates for the LMD of single tracks made of GE4822 of about 5 to 10 10⁴ K/s for beam diameters of 0.5 to 2 mm are simulated. The process conditions used in the study for the modeling differ significantly in the preheating temperature, which is about 400 °C, so that a reduction of cooling rates with increasing preheating temperature is to be expected.

Figure 8 shows the solidification rates in scanning direction at the melt pool boundary for the different processing parameters.

Solidification rates < 0 K/s indicate melting of the material. In comparison, solidification is faster for smaller beam diameters because of the fast scanning velocity used. Thus, a clear expression of dendritic microstructure morphology is expected in samples built with small beam diameters (0.6 and 1 mm). At the same time, the cooling rate for calculations with d_S = 0.6 mm is considerably higher than with 1 or 2 mm beam diameters, so that a comparatively finer microstructure is expected.

The respective maximum melt pool temperatures, maximum temperature gradients to the solid phase and maximum cooling rates are listed in Table II to allow quantitative comparison.

The results of the melt pool simulation using the measured preheating temperatures (see Figure 4), which depend on the build height, are shown in Figure 9. It is evident that there is no significant dependence of the cooling rate on the build height for d_s = 1 and d_s = 2 mm. This fits reasonably well with the small differences in the preheating temperatures. Therefore, a correspondingly homogeneous microstructure can be expected. The cooling rates when using a smaller beam diameter vary according to the local preheating temperatures over the build height, but are all in the range of 10³ K/s.

The influence of the process parameters used on the microstructure is expected to be visible. For the sample built with d_s = 0.6 mm, the microstructure is considered separately as a function of the build height because of the significant differences expected.

The macroscopic homogeneity of the samples can be examined exemplarily in Figure 10. Identical high individual layers are visible over the vertical structure with differences between the track and track edge areas that can be recognized by the constantly alternating shading.

In addition, Figure 10 also shows constant differences between the sample edges (left and right) and the middle of the sample. When looking at individual sections in detail, the process-related track and layer structure can be seen (Figure 11). Slight differences in microstructure fineness and morphology are visible depending on whether track edge or middle areas are considered. Beyond these regularly alternating areas, no differences in the microstructure are visible depending on the height of the structure.

After heat treatment, individual traces and layers are no longer visible, and the microstructure is homogenized (Figure 12).

Hardness measurements taken vertically in the middle of the as-built sample (Figure 13) show no significant, but nevertheless recognizable tendencies of hardness, being dependent on the build height. It was assumed that especially the fineness and thus the hardness of the microstructure could depend on the height of the structure because of the differences in preheating and reheating during the process. This assumption is confirmed by neither SEM images nor this measurement’s results.

The increased reheating due to the build-up of superimposed layers in the lower sample areas and the overall higher temperature nevertheless might result in differences in the local phase distribution and thus in lower hardness of the material reheated in this way over a longer time period. Additionally, a harder top zone due to a higher oxygen content through surface oxidation of the edge area is assumed to be the cause for the notable increased hardness values on the upper edge of the sample, even after heat treatment. While the samples slowly cool down after the process, the additional shielding gas flow is off, which allows for additional oxidation and oxygen diffusion in all near-surface regions. This is not essentially relevant for application, as surface regions will regularly be removed in subsequent finishing steps.

In Table III, the measured oxygen and aluminum contents of the examined samples are listed.

The oxygen content increases significantly as the beam diameter is reduced, although the material and process conditions are identical. The only other source of oxygen besides the powder material is the residual oxygen content of the inert gas atmosphere, which is <50 ppm O during the process. All of the process parameter-dependent influences during the process are complex, such as the cooling effect of inert gas flows, shading effects caused by powder particles and surface effects. However, the correlation between laser beam diameter or build-up rate and oxygen content is obvious. The reason for this is due to the smaller melt pool surface area, integrated over the process time, exposed to the residual oxygen-containing process atmosphere when using a larger laser beam diameter. The integrated melt pool surface A_S of a sample is the result of the calculation based on melt pool width b_S, track distance ∆y_s, layer thickness Δz and sample height H, width B and length Lwithout consideration of edge effects and surface curvature. Values for the process parameters used, in each case normalized by division by the sample volume, are listed in Table IV.

The cooling rate of the melt is regarded as a further critical variable. This determines the dwell time in the molten state during which oxygen absorption primarily takes place. At d_S = 2 mm, the cooling rate of the melt is about half as great as when using a laser beam diameter of 0.6 mm. Nevertheless, the total surface per time exposed to the environment in the molten state is still much smaller when using the larger beam diameter.

The global aluminum evaporation for all investigated samples is coherent with the literature references. This amounts to <1 wt pct evaporation for LMD and LPBF and approximately 1 wt pct or higher evaporation in EBM.[8,12,24] The evaporation is highest when using a beam diameter of 0.6 mm. According to Reference 24 evaporation is primarily proportional to the line energy E_S. Since evaporation is a surface effect, the beam diameter d_S is also relevant. Thus, the surface energy E_A can be formulated aswith laser power P_L and scan speed v in mm/s. This results in the values listed in Table V for the area energy of the parameter sets used.

It is assumed that during the process the aluminum evaporation temperature (2450 °C) is reached and, with greater surface energy, also higher temperatures at the surface of the melt pool. The use of the surface energy as the sole criterion is therefore not consistent with the measurement results, as this would imply the aluminum evaporation is highest when using a beam diameter of 2 mm. If the melt pool temperatures (Table II) are taken into account, a maximum aluminum evaporation would have to be assumed when using a beam diameter of 1 mm. Thus, the measurement results can be mainly justified by the larger melt pool surface that adds up over the total time of the build-up when using a beam diameter of 0.6 mm. In total, no definite statement can be made because of both the small deviations of the measured values from one another and the standard error for chemical analysis being expected to be in the range of at least 0.1 pct.

Figure 14 compares LMD microstructures of GE4822 observed using SEM-BSE with different beam diameters. The grain size increases visibly using a larger beam diameter (Figures 14(b) through (d)). Uniformly visible is the heterogeneous microstructure with very fine lamellar, non-measurable amounts of α₂ + γ as well as structureless γ matrix with blocky single α₂ phase.

The LMD process itself provides high cooling rates and numerous cycles of reheating at elevated temperatures by which the resulting share of lamellar α₂+ γ is supposed to be influenced. According to Reference 23 high cooling rates support the α → α₂+ γ transformation by extending the transformation range even resulting in feathery or Widmanstaetten features. This leads to the assumption that with higher cooling rates—in the study represented by using a beam diameter of 0.6 mm—the share of lamellar phase is higher compared to when using a larger beam diameter and thus provides lower cooling rates, respectively. However, the amount of γ seems to be especially high for the smallest beam diameter. For samples made of GE4822 produced by EBM, Reference 25 discusses to what extent the α phase is the result of incomplete massive γ transformation or nucleation from massive γ phase at high process temperatures. The literature indicates that cooling rates of several 100 K/s[26] are sufficient to trigger the massive γ transformation. In the AM process, the conditions for the formation of massive γ are then given. Reference 27 for example, shows how α₂ plates are formed within γ_m grains of a Ti-46Al-9Nb alloy by 1-hours heat treatment at 1200 °C. The process behind this is a transformation of γ_m into a α₂+ γ duplex structure. Regarding the α₂ phase visible in Figure 14(a), an interpretation of this is at least locally obvious, as can be seen from the “nested” α₂ within γ. In Reference 28 investigations on Ti-48Al-2Nb-2Cr (GE4822) show that a heat treatment at 900 °C to 1000 °C, held for 30 seconds, could be sufficient for the transformation of the α₂ plates. Due to this and the knowledge from Figure 3 that these temperature ranges are reached and passed through several times by already welded solid state material, the variant of α₂ nucleation from a massive γ phase discussed in Reference 25 seems probable. No indications of massive γ formation are visible in the microstructures of samples built up with d_S = 1 and 2 mm. This observation is consistent with the description documented in Reference 29 that larger proportions of massive γ phase are formed with higher cooling rates.

Figure 15 shows the microstructures of GE4822 material laser clad with different parameter sets after 4 hour of heat treatment at 1250 °C followed by cooling in air.

After the heat treatment, a microstructure consisting of α₂-Ti₃Al and γ-TiAl is present in all samples. In cross sections of the heat-treated samples built with 1 and 2 mm beam diameter, in contrast to the one built with 0.6 mm d_s, larger portions of lamellar γ + α₂ grains are visible. According to this, the phase transformation during cooling takes place via the γ + α₂ phase region. The α₂ phase fraction is visually smaller in all samples than before the heat treatment.

From the hardness test (Figure 16), despite the visibly coarser grains in samples built up with d_S = 1 and 2 mm, no significant difference is apparent. However, the visual impression that in the as-built state the proportion of softer y-phase (compared to the α₂ or lamellar phase) is higher in samples produced with d_s = 0.6 mm is coherent with the results. The lamellar distances are also an influencing factor to be considered.

With heat treatment, existing differences in the microstructure within the respective samples are almost completely compensated (Figure 12). The microhardness is reduced throughout, mainly by grain coarsening and the decrease of α₂ phase fraction. Within the scope of the measurement deviations, no significant differences were found even after heat treatment.