Enhanced Densification of PM Steels by Liquid Phase Sintering with Boron-Containing Master Alloy

Powder metallurgy (PM) steels for high-performance applications are of significant interest because of their low costs, flexibility in alloy design, and high volume production. However, reaching high density is not viable through conventional pressing and sintering which limits the usage of PM steel components in high-performance and demanding applications. Hence, achieving high density is essential to obtain the demanding mechanical properties that will enable PM parts to serve as an alternative to conventionally manufactured parts. High densities can be achieved either through pressure-based or sintering-based techniques or by combining both.[1‐4] Sintering with liquid phase forming additives activates the sintering and enhances the final densification through the liquid phase sintering mechanisms (LPS).[5,6] Elements with lower atomic number such as P, B, and C are well suited for activating sintering of iron-based powders due to the difference in electronic structure, phase composition, and grain boundary cohesion.[7] Lower amounts of such liquid-forming additives results in better densification by the liquid phase sintering (which is a dominant mechanism) than the surface energy activation.[5]

Boron addition significantly improves the density of the PM steels by forming a liquid with iron at ~1447 K (1174 °C) due to the eutectic reaction γFe+Fe₂B→L.[8] Once the liquid forms, more iron will be dissolved in the liquid, which results in persistent liquid formation. In PM steels, boron is introduced either as elemental boron[7,9‐14] or ferroboron,[15‐20] or master alloy,[21‐24] or compounds like hBN or B₄C.[25,26] Enhanced density levels were reached with elemental boron addition for both the carbonyl and water-atomized powders.[7] Ferroboron tends to agglomerate, resulting in inhomogeneous density distribution during sintering.[26] Selecká et al. [15] observed an increase in density by introducing ferroboron in Fe-Mo-prealloyed systems only when the boron content reaches beyond 0.2 wt pct. The design of master alloy is based on proper selection of alloy systems which can melt at lower melting temperatures so as to enable liquid phase sintering.[27] The elements such as Mn, Si, Cr, Ni can also be introduced without affecting the base powder compressibility.[28] When boron is introduced in the master alloy, it provides the flexibility of generating a liquid phase from the melting of the master alloy and the eutectic reaction between iron and boron.

In this study, the water-atomized iron powder (ASC100.29) and Fe-Mo-prealloyed powder (X-Astaloy 0.45Mo) from Höganäs AB were used as the base powder. Gas-atomized Ni-Mn-B master alloy (MA) powder, known as CEITALOY HD, developed by CEIT, Spain[22,29] with the nominal composition of 46 wt pct Mn, 46 wt pct Ni, and 8 wt pct B was used. The MA powder was produced by Höganäs AB using gas atomization and sieved below 45 µm. The MA powder was admixed with the base iron powders along with 0.4 wt pct Kenolube lubricant. The powder mixes were uniaxially pressed to impact bars with dimensions (10 × 10 × 55 mm³) according to ISO 5754 to reach ~7.3 g cm⁻³ density. The influence of carbon was studied by admixing 0.3 wt pct of natural graphite (UF4 from Kropfmühl). The master alloy addition of 2.5 and 1.5 wt pct was studied, which leads to an overall boron content of 0.2 and 0.12 wt pct, respectively. Higher MA addition was studied initially and then MA content was optimized to improve the mechanical properties of the components. Table I shows a summary of all the combination of powder mixtures and the designation that will be used henceforth. The six samples were chosen to study the influence of carbon, molybdenum, and MA content on densification behavior and mechanical properties.

All the samples were de-lubricated in a laboratory tube furnace at 723 K (450 °C) for 30 minutes under pure N₂ atmosphere. DSC was performed on de-lubricated samples with a heating and cooling rate of 10 K/min in high-purity argon (99.9999 pct) to study heat and phase changes. Simultaneous TG/DTA/DSC thermal analyzer STA 449 F1 Jupiter® from Netzsch was used for DSC measurements. All the laboratory sintering experiments were performed in a Netzsch 402L push rod dilatometer to detect the dimensional changes associated with sintering. Sintering was performed at 1513 K (1240 °C) for 30 minutes in Ar-50H₂ atmosphere with the heating and cooling rate of 10 and 30 K/min, respectively. In order to detect the formation and distribution of liquid phase during the heating stage as well as its effect on microstructure formation, interrupted sintering trails were performed at temperatures of 1273 K, 1373 K, and 1513 K (1000 °C, 1100 °C, and 1240 °C) for 1 minute with the same heating and cooling rates as above. All sintered samples were tempered in air at 473 K (200 °C) and then impact tested. Density measurements were performed by measuring the dimensions and sample mass for green compacts and by Archimedes’ principle according to ISO 3369 for sintered samples. For metallographic investigations, the samples were grounded, polished, and then etched using 3 pct nital. A Leica DMRX optical microscope was used for microstructural investigation. LEO 1550 Gemini scanning electron microscopy (SEM) combined with energy dispersive X-ray analysis (EDX, X-Max, Oxford Instruments) was used for microscopy study. The fractography study was performed on the fracture surfaces obtained after impact testing using SEM. Hardness was measured using Wolpert Dia Tester with Vickers scale of 10 kgf (HV10). Thermo-Calc software using TCFE8 database, based on the CALPHAD method,[30] was used to identify temperature and the volume fraction of liquid phase formed under equilibrium conditions.

The addition of boron-containing master alloy to iron and iron-molybdenum PM compacts reveals significant shrinkage during sintering above 1373 K (1100 °C). The sintering process and densification are activated due to the liquid phase formation. Sintering follows the mechanisms of liquid phase sintering,[5] starting from melting, spreading, rearrangement, and shrinkage. The results from the dilatometry, DSC, and microstructural study show two regimes of liquid phase formation: LP-1 and LP-2. The formation of initial liquid phase LP-1 is due to melting of the master alloy ~1292 K (1019 °C) from DSC analysis[29] and is in the same range as the liquidus temperature for the Ni-Mn binary system ~1293 K (1020 °C)[31] and the eutectic temperature of the Ni-B binary system ~1291 K (1018 °C).[32] The dilatometry curves from Figures 1(b) through (d) indicate that the shrinkage rate is rapid above ~1373 K (1100 °C). This is due to the initial rearrangement of the powder particles from the prior master alloy melting (LP-1). After that, the eutectic formation occurs (LP-2) as evident from the endothermic peaks of DSC, see Figure 2. This is also shown by Thermo-Calc simulations having similar liquid formation temperature, see Figure 3(a). Also, the LP-2 stage as indicated by the dL/dt of the dilatometry curves in Figures 1(b) through (d) can be related to the initiation of the eutectic melting, where the shrinkage rate starts to deviate.

This melting behavior is enormously effected by the presence of Fe and Fe-Mo as it depends on the thermodynamic stability of the boride phases formed with Ni and Mn. Mo in Fe is homogenously distributed within the iron powders as it is prealloyed, and the presence of B from master alloy favors the formation of molybdenum borides since the free energy of Mo₂B and MoB₂ is much lower than the corresponding Fe borides.[17] For both Fe and Fe-Mo systems, the eutectic formation occurs with the following reaction: γFe+Fe₂B→L, and γ(Fe, Mo) + (Fe, Mo)₂B→L. This eutectic formation will dissolve more Fe and generates more liquid phase resulting in the accommodation of grains, pore elimination, and grain coarsening. Sarasola et al. have reported that increase in the Mo content up to 3.5 wt pct suppressed the liquid formation, as all the B in the system is expended by Mo to form borides.[14] Hence, having a lower Mo/B ratio is an advantage as less Mo will be available to form borides. For Fe-Mo+2.5MA and Fe-Mo+1.5MA with and without carbon, the shrinkage initiation is at a similar temperature; however, the shrinkage is significantly larger for Fe-Mo+2.5MA as observed from the dilatometry curves, see Figures 1(c) and (d). From Tables II and III, the experimental results of the dilatometry curves and DSC traces, it is evident that Mo addition shifts the liquid formation stages LP-1 and LP-2 to lower temperatures, but the simulations from Thermo-Calc, see Figure 3(b), showed a higher LP-2 formation temperature. However, both experiments and simulations agree that Mo presence contributes to the formation of a higher amount of liquid volume fraction, see Table IV, but the delay in formation of LP-2 temperature from the predicted simulations is due to the equilibrium conditions.[9] Results clearly show that the amount of liquid phase and the melting temperatures are slightly lower as simulated by Thermo-Calc and this is related to some inaccuracy of the thermodynamic data for the system. In addition, during simulations in Thermo-Calc, a homogeneous system is assumed, which is not the case for the Fe-MA powder system, characterized by large concentration gradients between the sites of MA location and core of the base iron/steel particles.

Along with prealloyed Mo, carbon also has a significant effect on the densification which is evident from Figure 6 for the samples with and without carbon. Its effect is more pronounced when sintered above 1373 K (1100 °C), that is, during the second stage of the liquid formation (LP-2). Both the DSC results and Thermo-Calc simulations from Figure 3(a) show that the carbon presence lowers the LP-2 formation temperature as observed elsewhere.[9,21] This also results in higher liquid molar volume fraction at 1513 K (1240 °C) as indicated in Table IV, which enhances the density levels. Also, increasing master alloy from 1.5 to 2.5 wt pct facilitates the formation of more liquid volume fraction which results in an increase in densification from 2 to 5 pct depending on the amount of carbon content. Introducing boron as a master alloy has a significant impact on the densification, as observed from this study

Ni addition accelerates the boron diffusion in the γ-iron, similar to the carbon behavior in the presence of Ni[33,34] and increases the density of the PM steels.[15,35] Liu et al.[36] have shown near-full density after sintering pure iron with 1 wt pct Ni and 0.3 wt pct B. As indicated by Wu et al.,[37] the eutectic structure consists primarily of M₂B or M₃(B,C) with Fe, Mo, and Mn being the main elements with the presence of carbon and boron and Ni is absent from the eutectic structures as it will be distributed in the iron matrix. This can also be observed from the EDX analysis of the boride network as shown in Figure 7(b). The presence of Mn does not show much influence overall but is observed along with Fe in the boride network.

Sintering with boron addition forms eutectic phase as a continuous layer along the grain boundaries, especially for 2.5 MA samples (0.2 wt pct boron), even though adding more boron and carbon increases the volume fraction of liquid which eventually increases the density but not necessarily the mechanical properties as it results in embrittlement. The eutectic phases along the grain boundaries are extremely hard, which eventually results in the inferior mechanical properties see IE values in Table V and have brittle failure see Figure 5. However, the 1.5 wt pct master alloy addition (0.12 wt pct of boron) gives enhanced ductility with the impact energy (IE) values greater than 100 J (see Table V), which is similar to the case of Fe-0.1B.[19] This is due to the formation of a discontinuous eutectic network along the grain boundaries as seen from Figure 7(a), where the borides formed are evident from Figure 7(b). Hence, the design of this alloy system is to avoid the formation of continuous boride network and to substitute discontinuous boride networks along the grain boundaries, as seen from Figures 7(a) and (b). In this way, intergranular embrittlement can be fully avoided, as observed in the case of stainless steel based PM systems where fully dense and ductile components with elongation as high as ~ 34 pct are obtained for 316 L.[23]

In this study, the liquid phase sintering of PM compacts with master alloy containing boron (Ni-Mn-B) added to water-atomized Fe and Fe-Mo-prealloyed base powders was investigated. Analysis of the densification behavior and identification of the liquid-forming stages during sintering through dilatometry and DSC analysis were performed. Carbon addition and MA alloy content were varied to observe their effect on the liquid formation temperature, properties, and performance.

Boron control is essential to reach the sufficient densification and optimum mechanical properties which are suitable for the high-performance applications without compromising the dimensional stability of the PM components.