*3.1. Structural Morphology of Milled Powders*

Pre-alloyed mixtures of the powder particles consisting of ferro-alloys and traces of elemental powders that were blended inside a tungsten carbide vial are shown in Figure 3a. From the micrograph, it can be observed that the ferro-alloy particle mixtures consisted of various shapes such as irregular dentritic and flaky structures. Similarly, the morphology of the mixtures is shown in Figure 3b,c for a milling condition of 5 h and 10 h, respectively.

**Figure 3.** (**a**–**c**) Morphology of milled powders at regular milling interval.

From the SEM morphology, significant variations in size and morphology of the prealloyed powders during the entire milling time at regular intervals can be observed. The initial condition of the powder particles was soft and ductile. As the milling was carried out in wet conditions, these soft particles tended to agglomerate with each other during the early stages. With the continuation of milling, the agglomerated particles break down and flatten. It is evident that the milled powders are finer when compared to the initial condition after 5 h of milling. The formation of finer particles is due to the repeated fracturing and welding between the powder particles due to collisions between ball-powder-vial. Due to the high impact ball-powder-vial and shearing action between them, a high amount of energy was transferred to the milled powders. However, after 10 h of milling, the milled powders were found to be fragmented due to continuous work hardening. The surface properties revealed that the fracturing of the powder particles was dominant when compared to cold welding.

From Table 2, several studies carried out on the formation of austenitic stainless steel through MA are presented and it can be inferred that the formation of austenitic stainless steel depends on the milling parameters, namely starting materials, milling medium, BPR and milling time.


**Table 2.** Evolution of various austenitic grade stainless steels from different starting powders through MA.

When compared to stainless steel balls and vials, the kinetic energy of the moving particles within the vial is higher for tungsten carbide due to its higher density. If the austenitic stainless steel powders were to be developed from elemental powders, their evolution would require much more time compared to their development from pre-alloyed mixtures.

#### *3.2. Evaluation of Density for the Hot Pressed Samples*

Consolidation of the MA powders was carried out using hot pressing under a vacuum atmosphere. Density measurements for the vacuum hot pressed samples were performed using Archimedes principle. Theoretical density was calculated using rule of mixtures and compared with the hot pressed density as shown in Figure 4. The density attained for the hot pressed samples is averaged to be 7.67 g/cc which is 98.8% the density when compared to the theoretical density. A higher density was achieved due to MA, which fractured the powder particles to produce complicated shapes, which in turn increased the surface area of the powder particles required for hot pressing. Reduction in the size of the particles led to the formation of single compacts during hot pressing through grain boundary diffusion rather than lattice diffusion [29,30]. The high specific area of the powder particles initiated the surface diffusion mechanism during the initial stages of sintering. As the sintering progressed, grain boundary diffusion occurred and during later stages volume diffusion took place to produce highly dense compacts.

**Figure 4.** Density of hot pressed samples developed through powder metallurgy route.

#### *3.3. Metallography Analysis*

Figure 5a depicts the optical micrograph of 21-4N austenitic steel developed through vacuum hot pressing. The microstructure consists of equiaxed austenitic grains with grain boundary carbides distributed heterogeneously along the grain boundaries. The formation of grain boundary carbides is due to the slow cooling rate of 25 ◦C/min attained during vacuum hot pressing. The cooling rate plays an important role in the formation of metal carbides in an austenitic structure. This 21-4N alloy consists of low nickel and a high concentration of chromium in the metal matrix. Due to this, the grain boundary carbides consist of chromium-rich precipitates which are seen as dark phases within the micrograph. The light region represents austenitic grains. As the temperature for sintering is high (1200 ◦C), pores are not evident in the micrograph. Growth of grains did not occur at a faster pace at the beginning stages of sintering. However, at later stages of the sintering, cycle due to high temperatures, the available pores were rounded off. This can be interpreted in relation to a larger grain size in comparison to the crystallite size obtained as a result of MA. The sintering mechanism promotes the volume diffusion process through grain boundaries which act as vacancy sinks. The coalescence of pores occurs at high temperatures where grain growth is observed with grain boundaries as pinning points. Grain growth resulted in the rounding off of pores with little room for vacancies. Similarly, the density of the samples increased due to grain growth as well as volume diffusion.

**Figure 5.** Optical micrograph of (**a**) 21-4N, (**b**) corresponding grain size.

The mean grain size of the austenitic grain was found to be around 7–12 μm as shown in Figure 5b. The grain size was measured using the intercept method as per ASTM E112. The size of the grains was much smaller compared to a similar cast product which contained 16 μm and 17 μm sized grains as reported by Ji et al. and Kumar et al., respectively [20,31]. The initial starting size of the powders followed by the MA inhibited the grain growth during sintering in order to obtain a fine grain size.

The precipitation of carbides along the grain boundaries can still be authenticated using a SEM-EDS analysis. Figure 6a,b show the SEM-EDS analysis of both the grain and grain boundaries of 21-4N austenitic stainless steel. The presence of carbide precipitates is due to the effect of alloying elements within the metal matrix which change the solubility of carbon in the austenitic matrix. As nickel and manganese are austenitic stabilizers, they enhance the precipitation of carbides by reducing their solubility within the metal matrix.

**Figure 6.** SEM-EDS analysis of 21-4N both at (**a**) grains and (**b**) grain boundaries.

In both the grains and grain boundaries, three sections of EDS measurement were selected and shown in Table 3. A variation in the chemical composition can be seen from the EDS analysis. It is evident that elemental migration took place from the grains to the grain boundary. From Figure 6a, concentration spots (1, 2 and 3) within the grain reveal the presence of an austenite stabilizing element such as nickel at an appropriate level. Similarly, at the grain boundary spot in Figure 6b, a higher concentration (1, 2 and 3) of chromium, carbon and manganese along with iron were found in excess in comparison to the required elemental composition. In comparison to similar studies conducted by Xu et al., the size of the carbide precipitates were not larger or continuously dispersed within the metal matrix [32]. Moreover, the development of the alloy through the powder metallurgy route led to the formation of fine grains and restricted the size of metal carbide precipitates.



The XRD analysis of the vacuum hot pressed steel is shown in Figure 7, and it can be confirmed that these metal carbides were mostly composed of M23C6 where M represents chromium, manganese and iron. The formation of carbides occurred mainly due to the higher C/Cr ratio. From the metallographic studies, it can be understood that M23C6 nucleates very easily at the beginning when compared to other types of carbides. Moreover, the structural shape of these metal carbides is of a globular or cellular type. These carbide precipitates are usually observed near the grain boundaries due to it being a favorable place for carbon, with an interstitial solid solution capability which rapidly diffuses near the grain boundaries. The intensity of austenite peaks was dominant when compared to M23C6 peaks. The formation of the M23C6 peak in hot pressed samples is due to slow cooling in the sintering cycle. During the sintering of MA powders at 1200 ◦C, the peaks of austenite became intense and sharper with diminished peaks of M23C6. The presence of M23C6 at the grain boundary improved the material resistance to heat and sliding of grains. Similarly, Figure 8 shows the XRD analysis of powder particles milled for 10 h. The austenitic peak is evident from the XRD analysis.

**Figure 7.** X-ray diffraction pattern for vacuum hot pressed sample.

**Figure 8.** X-ray diffraction pattern for milled powder.

Figure 9a represents the TEM microstructure of the vacuum hot pressed steel. Within the austenitic grain, twins were available which indicates the presence of an austenitic matrix. Face centered cubic crystal alloys have a tendency to form twins within the microstructure when subjected to cold working and annealing as a result of low stacking fault energy. Due to low stacking faults, the energy required for twin formation is lower when compared to grain boundary formation. These kinds of fine austenitic twins are found as the material is subjected to cold working and annealing during ball milling and vacuum hot pressing, respectively. Few dislocations were also found within the grains due to hot pressing. The formation of nano crystalline grains along with dislocations is the outcome of severe plastic deformation. The dislocations are pinned at the grain boundaries. This is due to the presence of metal carbides which are present at the grain boundaries that resist dislocation due to the precipitate strengthening mechanism. These precipitates will impede movement or dislocation and resists plastic deformation. Due to MA, the powders consists of residual stresses which act as sites for nucleation [33]. During sintering, these powders nucleates and give rise to the particle stimulated nucleation phenomenon. Face centered cubic structures are found to have particle stimulated nucleation. As the holding time during sintering is 2 h, some grains nucleate near the grain boundaries. This is because a favorable place for grain nucleation is near the grain boundaries which are also the locations for rapid diffusion when compared to a lattice. From Figure 9a, few grains are seen in between the grain and grain boundary region. Grain sizes of the TEM samples were determined by the method of linear intercept. The mean grain size of 1 μm was obtained for the samples developed through vacuum hot pressing as shown in Figure 9b. Cast products with similar compositions had large precipitates of carbides along with σ-phases which were not visible near the grain boundary in the current study [3]. Due to the particle stimulated nucleation phenomenon, very few residual carbides were seen within the metal matrix along with the formation of grains and grain boundaries. Similarly, as the material was manufactured using MA and subsequently consolidated using hot pressing, σ-phases and other intermetallics were not visible as sintering took place under a vacuum atmosphere. The selected area diffraction (SAD) pattern, as shown in Figure 9c, indicates the formation of a nanocrystalline structure even after the consolidation of MA powders at 1200 ◦C. The rings represent the formation of nano crystallite structures of the metal matrix. Correspondingly, the outer ring represents the formation of face centered cubic structures, namely (111) and (222), respectively. Both the austenitic matrix and precipitates of the carbides co-exist within the metal matrix due to the fact that both have face centered cubic structures.

**Figure 9.** TEM micrograph of heat resistant austenitic steel of 21-4N (**a**) bright field image where arrow mark indicates dislocation lines, (**b**) distribution of grain size for the austenitic matrix, (**c**) corresponding SAD pattern.
