Effects of Rare Earth Lanthanum on the Solidification Structure and Hot Ductility of Fe-43Ni Expansion Alloy

Comparison of macrostructures

The macrostructures of the tested cylindrical ingots are shown in Figure 2. The corresponding relationship of lanthanum contents and characteristic parameters of the macrostructures are shown in Figure 3. As clearly seen from Figures 2 and 3, the macrostructures of the tested alloys can be significantly refined with increasing lanthanum content in the range of 0~0.025 %, resulting in the reduction of the average equiaxed grain size and the improvement of the proportion of the equiaxed grains at the transverse cross-section; moreover, the macrostructure has completely changed into the equiaxed grains, especially with the addition of 0.025 % lanthanum. However, when the lanthanum content is up to 0.04 %, the macrostructure has started to become coarse again. Therefore, it can be concluded that the macrostructure of the alloys can be remarkably refined by adding 0.01~0.025 % lanthanum; with 0.04 % lanthanum, the refinement result was worse instead.

Figure 2:

Macrostructures of the tested cylindrical ingots (a) 0 % La, (b) 0.1 % La, (c) 0.025 La, (d) 0.4% La.

Figure 3:

Relationship between lanthanum contents and characteristic parameters of the macrostructures.

Hot ductility evaluation

The ingot of alloy 4 with addition of 0.04 % lanthanum cracked during the hot forged process, and this revealed the fact that the hot ductility of alloy 4 was very poor, as shown in Figure 4. The RA of alloys 1, 2 and 3 at 850–1200 °C is shown in Figure 5. According to Mintz et . [12], the temperature range in which the RA is less than or equal to 60 % is a crack sensitive range, which is called the hot brittle range. In terms of this rule, the alloy 1 has a good hot ductility in the temperature range 1100–1200 °C and the RA is above 70 %, which illustrates that the alloy is suitable for hot rolling provided the finish rolling temperature did not fall below 1100 °C. However, the RA is less than 50 % and the hot ductility is deteriorated when the temperature is between 850–1050 °C. The alloys 2 and 3 both have a good hot ductility in the temperature range 1050–1200 °C and the RA is above 70 %, moreover, the hot ductility is still good at 1000 °C in alloy 3 rather than alloy 2. Both alloys 2 and 3 have a poor hot ductility at 850–950 °C. Therefore, it can be concluded that lanthanum has a beneficial influence on the hot ductility of Fe-43Ni alloy. The addition of 0.01–0.025 % lanthanum can improve the hot ductility over the whole testing temperature range, especially at 1000–1050 °C.

Figure 4:

The cracked ingot during hot forged process.

Figure 5:

Hot ductility curves of tested alloys.

Fracture morphology

The fracture surfaces of the tensile-tested specimens of alloys 1 and 2 at 1000~1100 °C obtained by SEM are shown in Figure 6. At 1000 °C, the tensile-tested fracture of alloy 1 exhibits the intergranular brittle fracture, as shown in Figure 6(a), which agrees with the low ductility values given in Figure 5; however, alloy 2 shows a mixture of intergranular and ductile fracture, as shown in Figure 6(b). This result can indicate that the grain boundary strength of alloy 2 is higher than that of the alloy 1. With temperature increasing to 1050 °C, the fracture of the alloy 1 changes to the mixture of intergranular and ductile fracture, as shown in Figure 6(c), but the alloy 2 exhibits completely ductile fractures with large and deep dimples on the fracture, as shown in Figure 6(d). At 1100 °C, both alloy 1 and alloy 2 show completely ductile fractures, as shown in Figures 6(e) and 6(f).

Figure 6:

Fracture morphology of tensile-tested samples: (a) alloy 1 at 1000 °C, (b) alloy 2 at 1000 °C, (c) alloy 1 at 1050 °C, (d) alloy 2 at 1050 °C, (e) alloy 1 at 1100 °C, (f) alloy 2 at 1100 °C.

Grain refinement mechanism

It has been reported that the changes of solidification macrostructures are comprehensively competitive action between the nucleation and growth reactions of columnar and equiaxed grains, which in turn are strongly effected by the effective inclusions [13]. The effective inclusions inducing grains nucleation have a great impact on the solidification structural development, and if the effective inclusions were too few, the refinement of solidification macrostructures would be difficult. Compared with the above experimental results, it can be concluded that lanthanum containing inclusions may serve as the effective inclusions sites promoting the refinement of solidification macrostructures. Therefore, the category, size and quantity of lanthanum containing inclusions has been discussed in the following paragraphs.

The morphology and energy spectrum analysis of inclusions in tested samples treated with various lanthanum were analyzed using SEM and EDS. The typical inclusions are shown in Figure 7. The inclusions in alloy 1 mainly are Al₂O₃, the shape of which is irregular and the size is approximate 4 um, as shown in Figure 7(a). After adding 0.01 % lanthanum, the inclusion in alloy 2 are complete La₂O₂S, the shape of which changes into globular, and the size reduces to 2 µm, as shown in Figure 7(b). When lanthanum content is up to 0.025 %, the inclusions are still La₂O₂S, which are finer and spherical, distribute dispersedly, as can be seen in Figure 7(c). However, with increasing lanthanum content up to 0.04 %, the size of La₂O₂S inclusions grows up to 5 µm due to agglomeration of La₂O₂S inclusions, and the shape of which becomes irregular again, as can be seen in Figure 7(d).

Figure 7:

Morphology and energy spectrum analysis of typical inclusions in tested alloys.

According to Turnbull and Vonnegut [14], the inclusions could serve as the effective inclusions sites promoting the refinement of solidification macrostructures, the two following conditions are necessary: (a) the high-melting point of inclusions is higher than that of the matrix metal, (b) the lattice misfit of low-index surfaces between the inclusions and the matrix metal should be less than 12 %, and the less the lattice misfit, the more effective the inclusions serve as the effective inclusions sites.

The melting point of La₂O₂S and the Fe-43Ni alloy are approximate 1940 °C and 1400 °C respectively, thus it can be concluded that La₂O₂S is a high-melting point inclusion. The crystallographic data of La₂O₂S and the γ-matrix of Fe-43Ni alloy are shown in Table 2 [15, 16].

According to Bramfitt [17], the two-dimensional lattice misfit mathematical model is as follows:

where δ is the lattice misfit of the two interfaces, the (hkl)_s is the low index plane of the substrate of the inclusions, the [uvw]_s is the low index direction in (hkl)_s, the (hkl)_n is the low index plane in the matrix metal, the [uvw]_n is the low index direction in (hkl)_n, d_[uvw]s and d_[uvw]n are the atomic spacing along [uvw]_s and [uvw]_n, the θ is the angle between [uvw]_s and [uvw]_n.

Based on analysis above, the lattice misfit between La₂O₂S and the γ-matrix has been calculated, the results are listed in Table 3. It can be seen that two-dimensional lattice misfit between the face (0001) of La₂O₂S and the face (100) of γ-matrix is 5.42 %. Therefore, La₂O₂S inclusions can serve as the effective inclusions sites promoting the refinement of solidification macrostructures.

As mentioned above, the macrostructures of the tested alloys were first refined and then became coarse with increasing lanthanum content. That is because that the changes of the macrostructures relates not only to the category but also to the size and number density of the effective inclusions sites [18, 19]. To find out the quantity density and the size distribution of La₂O₂S inclusions in alloys 2, 3 and 4, the samples were mechanically ground, polished, etched and examined using SEM and EDS. The quantity of La₂O₂S inclusions in one square millimeter in each sample were calculated, and the results are shown in Figure 8. It is clear that the quantity density of La₂O₂S inclusions in alloy 3 is the largest of all, the quantity density in alloy 2 is between that of alloy 3 and 4, and the quantity density in sample 4 is the lowest. Furthermore, the size distribution of La₂O₂S inclusions in each sample also was counted. The results are shown in Figure 9. It is clear that the size distribution of La₂O₂S inclusions in alloy 2 are close to that in alloy 3, and the size are mainly concentrate on the range of 2~4 µm, but the size of La₂O₂S inclusions in alloy 3 is a little smaller than that in alloy 2. As to alloy 4, the size distribution of La₂O₂S inclusions are mainly focused on the range of 4~6 µm, which are larger than the two former. It has been reported that the most effective nucleation size for inclusions is about 2 µm during solidification process [20]. Not only the quantity density of La₂O₂S inclusions is the largest, but the size distribution of La₂O₂S inclusions also is the most appropriate in alloy 3. These statistical results are consistent with the results of microstructures in alloy 2, 3 and 4. Therefore, it can be concluded that the change of the quantity density and the size distribution of La₂O₂S inclusions are the main reason for the changes of the macrostructures.

Figure 8:

Quantity density of the effective La₂O₂S inclusions.

Figure 9:

Size distribution of the La₂O₂S inclusions.

Improvement mechanism of the hot ductility

The alloys 2 and 3 have a better hot ductility compared to the alloy 1 over the whole testing temperature range, especially at 1000–1050 °C. If the high temperature conditions were such that the grain boundaries are strengthened and the movement of them become possible during straining, the hot ductility would be improved [1]. Several structural features of the longitudinal section near the fracture may be responsible for the changes of the hot ductility, as shown in Figure 10.

Figure 10:

Macrographs of longitudinal sections close to the fracture of alloy 1: (a) alloy 1 at 1050 °C, (b) alloy 1 at 1100 °C, (c) alloy 2 at 1000 °C, (d) alloy 2 at 1050 °C, (c) alloy 3 at 950 °C, (d) alloy 3 at 1000 °C.

Figure 10(a) and (b) show the optical macrostructure of the alloys 1 after fracture at 1050 °C and 1100 °C, which are the crucial temperatures in the RA curves and represent the demarcation point of hot ductility of alloy 1. At 1050 °C, the equiaxed grains keep the original shape and the grain boundary separation is much more conspicuous, with large elongated macrocracks forming along the grain boundaries. These indicate that the grain boundaries are so weak that they separate before the grains experience any deformation, which are also the cause of bad hot ductility at 850–1050 °C. However, the optical macrostructure exhibits completely different structure changes at 1100 °C, the tip of the fracture are consisted of much finer equiaxed grains, suggesting remarkable evidence of dynamic recrystallization (DRX). It is reported that the occurrence of dynamic recrystallization has a substantial contribution to the high hot ductility [21]. Dynamic recrystallization, (i. e. grain boundary migration) has moved grain boundaries away from the voids and microcracks, led to isolation of microcracks and prevented their coalescence, resulting a good hot ductility above 1050 °C [22, 23].

Figure 10(c) and (d) shows the optical macrostructure of the alloys 2 after fracture at 1000 °C and 1050 °C. At 1000 °C, it can been seen that the grain boundary separation is not obvious, moreover, the grains of the fracture tip are elongated and exhibit a little plastic deformation, illustrating that the grain boundaries are strengthened and produce a little movement compared to alloy 1. At 1050 °C, the optical macrostructure of the fracture tip is the same with that in Figure 10(b) and exhibits complete dynamic recrystallization structure. Figure 10(e) and (f) show the optical macrostructure of the alloys 3 after fracture at 950 °C and 1050 °C. At 950 °C, the grain boundary separation is also not obvious and the grains of the fracture tip has the elongated tendency. At 1000 °C, the optical macrostructure of the fracture tip exhibits complete dynamic recrystallization structure.

In addition, the grains size of the tensile fractures directly represents the grain structure in hot forged condition, thus it can be conclude that the grain structure in hot forged condition can be gradually refined with increasing lanthanum content in the range of 0~0.025 %, which is the consistent with the macrostructures of the tested cylindrical ingots.

According to the experimental analysis above, it can been seen that the onset temperature for dynamic recrystallization is brought forward by 50 °C and 100 °C in alloys 2 and 3, respectively. Dynamic recrystallization needs thermal activation and initiates almost always at grain boundaries, it would accelerate at high temperatures or in the much finer grained structure condition [24]. The grain sizes of alloy 1, 2 and 3 in hot forged condition are gradually refined with increasing lanthanum content, which provides more initiation location for dynamic recrystallization, therefore, the onset temperature for dynamic recrystallization will be brought forward with the same thermal activation and strain rate during the hot tensile test. In addition, the quantity density of La₂O₂S inclusions also gradually increased with increasing lanthanum content which indicate that more La₂O₂S inclusions will locate at the boundaries. La₂O₂S inclusions are beneficial for the nucleation of favor for the γ-matrix and promote the dynamic recrystallization. Therefore, it can be concluded that the occurrence of dynamic recrystallization result in the improvement of hot ductility in alloy 2 at 1050~1200 °C and in alloy 3 at 1000~1200 °C.

As mentioned above, the grains of the fracture tip in alloy 2 and 3 are elongated and exhibit a little plastic deformation at no dynamic recrystallization temperature zone. Because the grain sizes of alloy 1, 2 and 3 are gradually refined with increasing lanthanum content, which increases more triple junctions of grain boundaries and the grain boundary densities. During hot tensile process, once the microcrack initiated at the grain boundary, though it may be able to penetrate a little distance into the grains, it would soon meet a new triple junctions of the boundaries and their direction change, moreover the penetration distance in the refiner grains is shorter than that in coarse grains, as shown in the Figure 11, which will restrict the growth rate of the macrocracks and strengthen the grain boundaries. Lanthanum is easy to segregate at the grain boundaries because the free energy of lanthanum atoms in the intragranular is higher than that at grain boundaries and the segregation of lanthanum at grain boundaries is favorable for the hot ductility. Sulfur is also easy to segregate at the grain boundaries even though with very low content and it is harmful to the hot ductility. The segregation of lanthanum at the grain boundaries decreases the segregation of sulfur at the grain boundaries, thus strengthening the grain boundaries. In addition, La₂O₂S inclusions capture part of sulfur in the grains and reduce the segregation of sulfur at the grain boundaries, so as to strengthen the grain boundaries. The strengthened grain boundaries have a more resistant to the grain boundary separation during hot tensile process, so the grains could experience plastic deformation, leading to the improvement of hot ductility.

Figure 11:

Schematic diagram of the microcrack penetration.

Deterioration mechanism of the hot ductility

The ingot of alloy 4 with addition of 0.04 % lanthanum cracked during the hot forged process, thus the hot ductility of the ingot was very poor. To find out the cause of deterioration mechanism, the samples from cracked parts were mechanically ground, polished, etched and examined using SEM and EDS. There are Fe-Ni-La intermetallic phases linearly dispersing along grain boundaries, as shown in Figure 12. It is known that adding appropriate rare earth can purify the molten steel, modify inclusions and improve the hot ductility [25]. However, if the lanthanum content was excessive for the matrix of the alloy, such as adding 0.04 % in this study, redundant lanthanum would segregate at the boundaries and react with Fe and Ni to form Fe-Ni-La intermetallic phases in addition to form La₂O₂S inclusions. The Fe-Ni-La intermetallic compound is a brittle phase, which not only weakens the grain boundary binding force, but also hinders the movement of the grain boundaries in the hot forged process, thud leading to the deteriorated hot ductility.

Figure 12:

Morphology and energy spectrum analysis of intermetallic compounds at grain boundaries.