Nanostructured materials have been shown to exhibit outstanding high strength and hardness, corrosion resistance, good soft magnetic performance, and other unusual properties 12 34 5. For establishing meaningful correlations between the structure and properties, it is necessary to have fully dense, artifact-free, and well-characterized samples. A number of prepared methods of bulk nanostructured materials have been investigated in the recent decades 678 9 10. These processes usually use a two-step approach including the synthesis of nano powder and particles and their subsequent consolidation by hot isostatic pressing, shock compression, powder sintering, etc, or a one-step approach such as electrodeposition, devitrification of amorphous and severe plastic deformation. In the two-step approach, the materials often suffer from porosity, contamination, and weak bonding resulted to low ductility, whereas the one-step approaches are limited to the synthesis of small-scale products. Therefore, it is of interest to develop new processes that are convenient, low in cost, and capable of being scaled up for tailoring the nanostructured materials, for instance a self-propagating high temperature synthesis (SHS) technique 11 1213.

The Fe-B alloy system is of interest in several fields of engineering materials. This system shows a significant glass forming ability (GFA) for compositions close to the Fe-rich eutectic 14 15. Nanostructured Fe-B alloy is usually obtained from the metallic materials 16 17. The stable phase Fe₂B and metastable phase Fe₃B often coexist in the nanostructured Fe-B alloy. The thermal stability of the nanostructured Fe-B alloy is reduced greatly because the metastable phase Fe₃B transforms into the equilibrium phase mixture of α-Fe and Fe₂B at 840 °C 18. In this paper, we have successfully synthesized bulk nanostructured Fe_94.3B_5.7 alloy making use of the one-step approach of an SHS combining a rapid cooling technique that is convenient, low in cost, and capable of being scaled up for tailoring bulk nanostructured materials 11121319. The nanostructured Fe_94.3B_5.7 alloy is only composed of α-Fe(B) and t-Fe₂B phases. It possesses good soft magnetic performance and excellent mechanical behavior.

The aluminothermic reaction was designed as expressed in Eq. 1. The reactant of ferric sesquioxide, aluminum, and boron powders were dry-mixed for 8 h in a hard steel vial with Al₂O₃ spheres, and then about 80 g of the mixed powders were cold-pressed in a copper mold (70 mm in the diameter) under a uniaxial pressure of 50 MPa. The characteristics of the reactant powders are given in Table 1. Three grams of Al, S, and MnO₂ powder mixture in the mass ratio of 1:1:1 was pressed into a pellet with dimensions of 15 mm (diameter) × 4 mm. The pellet was put on top of the pressed reactant powders as an igniter.

The copper mold with the reactants was placed in an SHS reactor as described in the literature 19 20 21. The reactor was purged with argon gas at room temperature, heated to 180 °C, and then purged for a second time. Heating of the reactor continued after introducing of 7 MPa of argon gas. The reaction of the igniter was started when the reactor reached 260 °C, which results in an instantaneous release of a large amount of heat that ignites the aluminothermic reaction and the synthesis reaction was subsequently finished in a few seconds. The products were kept in the reactor under argon gas pressure to cool down.

After cooling to room temperature, the products were taken out of the reactor. The resulting Fe_94.3B_5.7alloy was about 5 mm thick and 30 mm in diameter. The black Al₂O₃product was on top of the target product, which separated naturally from the desired Fe_94.3B_5.7alloy and could be removed by hand. The Fe_94.3B_5.7alloys prepared in the repeated experiments have the same phases, nanostructure, and properties as those mentioned below, which indicates that the processing is stable and can be conveniently reproduced.

The polished cross sections were investigated with X-ray diffractometry (XRD, Philips X’per) using CuKα radiation. The samples before and after compressive test were observed with a JSM-5600LV scanning electron microscope (SEM). Several specimens punched from the cross sections of the products were electrochemically thinned in an electrolyte of nine parts of methanol and one part of perchloric acid with twin-jet electropolish at −30 °C and were examined using a JEM-1200EX transmission electron microscope (TEM) with a magnetically shielded object lens pole-piece operated at 120 kV. The differential scanning calorimeter (DSC, STA 449 C) was employed to evaluate the phase transformations in Ar atmosphere with the heating and cooling rates of 20 °C/min. The saturation magnetization and coercive forces were tested using a vibrating sample magnetometer (VSM, 7304, USA) equipped with a 1.0 T magnet and the measurements were conducted at room temperature. The cylindrical compressive specimens with a dimension of Ø 3 mm × 4.5 mm were cut using an electrodischarging machine, and the specimen surfaces were polished with a 1,000-grit emery paper. A quasi-static uniaxial compressive test was performed at room temperature with a crosshead speed of 6 × 10⁻⁴ s⁻¹. In order to minimize friction effects, the specimen–die interfaces were lubricated with graphite.

The chemical reaction of the igniter is initiated at about 260 °C in the SHS reactor and instantaneously releases large heat, which ignites the reaction of the reactants. Then combustion wave of the reaction propagates from top to bottom of the reactant compact, and the reactants transform to the Fe-B alloy and Al₂O₃ where the combustion wave has passed. The adiabatic temperature (T _ad) of the combustion reaction (1) at 260 °C is calculated to be approximately 3400 °C 20 21.

The XRD pattern of the Fe_94.3B_5.7alloy is shown in Fig. 1. The Fe_94.3B_5.7alloy displays the composition oft-Fe₂B and α-Fe phases; however, the metastable Fe₃B phase is not found. The strongest intensity ratio of the (110)_α/(002)_tof the Fe_94.3B_5.7alloy indicates that the predominant phase identified is the α-Fe phase. It is in accord with the result of the calculated phase diagram.

Figure 2shows the SEM secondary electronic images of the Fe_94.3B_5.7alloy. The Fe_94.3B_5.7alloy is a composite structure of micrometer-sized dendrites dispersed in a eutectic matrix. The dendrite trunk thickness is about 5 μm, and the arm spacing of the dendrite is a few micrometers. The dendrite phase has tetragonal structure according to the phase identification by the selected area electron diffraction (SAED) analysis. The eutectic colonies distribute uniformly and the size is in the range 10–20 μm, and the content of the eutectic structure in the Fe_94.3B_5.7alloy is about 70%.

The eutectic matrix microstructure cannot be resolved using SEM. The TEM observations of the Fe_94.3B_5.7alloy are represented in Fig. 3. The nanostructured eutectic is found relatively disorganized (Fig. 3a). The TEM observations demonstrate that fine α/tphases inside the colonies are arranged in an alternating fashion. Figure 3b shows the high magnification of the zone I in the Fig. 3a. The corresponding SAED of the I zone (Fig. 3c) and the II zone (Fig. 3d) results confirm the existence of the two phases of the eutectic, α-Fe(B) andt-Fe₂B. Figure 3c also indicates that the precipitationt-Fe₂B is found in the α-Fe(B) phase.

Eutectic microstructures involve the simultaneous growth of two solid phases from the melt, which results from the cooperative growth of two crystalline phases by a discontinuous reaction. In this kind of process, there is no concentration difference across the reaction front, but short-range diffusion takes place paralleling to the reaction front and the two components separate into two different phases. In a very high super-heating liquid state about 3400 °C, the contaminants introduced from igniter are dissolved and purified in the metallic melt. Thus, the undercooling degree is greatly increased because there are no heterogeneous nucleation sites for the liquid to crystallize on. Crystal growth velocities of Fe and Fe₂B phases are a function of undercooling degree in the Fe-B binary alloy melt 24 . The growth velocity of the Fe₂B phase is always higher than that of the Fe phase, as boron content is below 15 at% in the Fe-B alloy under the large undercooling degree 22 23. So the effect of solute enrichment at the solid–liquid interface tends to increase the tip radius of the primary Fe₂B. This process will be persisted until the occurrence of stability of solute dendrites. Thus the dendrite phase has the t-Fe₂B structure rather than α-Fe(B) in the Fe_94.3B_5.7 alloy under the deeply super-cooling degree. The higher growth velocity of the Fe₂B phase also results in the breaking of the symmetry of lamella eutectic and leads to the formation of a hypereutectic structure (Fe₂B + eutectic). At the same time, the fine eutectic structure is formed because of the large degree of undercooling.

By the TEM analysis, it is found that solid/solid transformation takes place during the sample cooling. Taking the Fe–B binary phase diagram 24 into account, the γ-Fe(B) phase undergoes two solid/solid transformations. The formation of polycrystalline phase instead of a single crystal phase is attributed to the type of transformation. After the eutectic transformation, crystals of α-Fe(B) phase are nucleated within the γ-Fe(B) phase with decreases of the temperature 25 . The eutectoid transformation undergoes a similar process at 912 °C. As the temperature, is decreased the B solid solution degree in the Fe is reduced, so the Fe₂B grains are precipitated in the α-Fe(B) phase (Fig. 3b).

In order to understand in detail the overall phase transformation process (i.e., solidification and solid-state transformation), complementary DSC analysis is performed for the Fe_94.3B_5.7alloy. Typical DSC curves are shown in Fig. 4. For heating (curve (a)), three sharp endothermic peaks (at about 1135, 1150, and 1212 °C) are observed. The first endothermic peak is associated with the solid solution transform temperatureT _ssdetected at 1135 °C in the solid state, and the second one with the eutectic melt temperature at 1150 °C. On further heating Fe₂B starts to melt at 1212 °C, resulting in the third peak. For cooling (curve (b)), three exothermic peaks (1134, 1171, and 1198 °C) occur as well, also corresponding to theT _ss, eutectic melt and Fe₂B melt temperatures, respectively.

Figure 5 displays that the coercive force (Hc) value of the Fe_94.3B_5.7 alloy is approximately 11 A/m and saturation magnetization (Ms) is up to 1.74 T. It is well known that the microstructure essentially determines the hysteresis loop of a ferromagnetic material 26 27. The low Hc is attributed to the finest composite structure. The lamellar spacing of the eutectic is below 50 nm, and the interval of the precipitation grain is down to 20 nm. These values are less than the ferromagnetic change length, which is typically 40–50 nm in the Fe-B alloy. Thus, the lower coercivity is obtained. Also, we can see that the high saturation magnetization reaches 1.74 tesla because the content of Fe element reaches 94.3 wt% in the Fe-B alloy.

Figure 6 shows the compressive engineering stress–strain curve of the Fe_94.3B_5.7 alloy. The yield strength of 1430 MPa, and impressive ultimate strength of 1710 MPa and plastic strain of 19.8% are reached in compression, respectively. The yield strength is far higher than the reported 580 MPa of the coarse crystalline Fe₈₃B₁₇ alloy prepared by the traditional technique, and the plasticity strain retains a high value at the same time 28 . The Fe_94.3B_5.7 alloy presents a continuous stress increase with increasing strain, pointing to continuous work hardening until failure.

Figure 7a presents the typical features on the fracture surface of the Fe_94.3B_5.7 alloy, illustrating a visible shear fracture, which is usually generated by plastic deformation. The remelting evidence and flat fracture surface with distinct viscous shear flow traces are observed in Fig. 7b, indicating that shear bands induce a severe local energy transformation during the fracture of the sample, resulting in local temperature increase and local ‘softening’ of the alloy at the moment of fracture. The Fig. 7c and d demonstrates an obviously viscous shear flow phenomenon. Some shear bands propagate round the dendrite, and others are arrested near the intersection of the eutectic and dendrites. The eutectic colonies around the dendrites markedly rotate to accommodate the local stress deformation, which avoids catastrophic fracture due to the local stress concentration 29 30. The composite microstructure of the Fe_94.3B_5.7 alloy is responsible for continuous work hardening in the compressive deformation process, which results in the large plasticity.

In this work, we applied a simple and effective means (SHS Combining Rapid Cooling Technique) to prepare the bulk nanostructured hypoeutectic Fe_94.3B_5.7alloy. The Fe_94.3B_5.7alloy is characterized by a lamellar eutectic witht-Fe₂B and α-Fe phases. The eutectic laminar space is about 50 nm and the size of spherical eutectic colonies is in the range 3–25 μm. The nanostructured eutectic is attributed to deep undercooling. The Fe_94.3B_5.7alloy displays good soft magnetic behavior with the coercive force (Hc) of 1.9 A/m and saturation flux density (Bs) of 1.75 T. The Fe_94.3B_5.7alloy also exhibits simultaneously high yield strength (1430 MPa) and large plasticity (~19.8%), which is attributed to the nanostructured eutectic with a few micrometer dendrites of the composite Fe_94.3B_5.7alloy. The shear bands induced in the eutectic matrix propagated is rounded or arrested by the dendrites. The SHS technique combined with rapid cooling technique can be used for tailoring nanostructured composite materials, which are expected to be applied to many sample alloy systems due to their extraordinary properties.