2020 Volume 61 Issue 3 Pages 548-556
TiB2 reinforced Fe matrix composites were investigated for their potential as a new generation of hot work tools which are mainly characterized by high thermal conductivity and high hardness in comparison with conventional materials. In this work, Fe–30 vol%TiB2 composites were sintered at 1373 K for different holding times (0, 0.3, 0.6, 1.8 and 3.6 ks). Apart from Fe and TiB2, newly formed phases of Fe2B and TiC were found in all sintered compacts. A good Fe/TiB2 interfacial cohesion was confirmed at atomic level at 0 ks, which was due to the occurrence of the special orientation relationship between {110} planes of Fe and $\{ 10\bar{1}0\} $ planes of TiB2. The observation of dislocations in TiB2 particles, attributed to the activation of slip systems, showed the plastic deformation ability of TiB2 at high temperature. The reaction between Fe and TiB2 was due to TiB2 dissolution in Fe at 1373 K and different diffusion depth of B and Ti atoms in Fe. Consequently, B directly reacted with Fe, since the solubility of B atoms was low in both α-Fe and γ-Fe. TiC probably precipitated from Fe–Ti–C solid solution along Fe grain boundaries in the cooling stage after sparking sintering, leading to a layer of Fe wrapping around TiB2. Among all the compacts, the one sintered at 1373 K for 0.6 ks displayed the excellent properties which were comparable in Vickers hardness and 133% higher in thermal conductivity, compared with that of SKD61 as the commonly used practical material. This work provides a new perspective to fabricate a future generation of hot work tools.
Fig. 5 HAADF images of Fe–30 vol%TiB2 compacts sintered at (a) 1373 K for 0 ks and (b) the high magnification image corresponding to the area in (a) marked by dashed box; (c), (d), (e) and (f) placed on the right are mapping analysis images of Fe, Ti, B and C corresponding to (a).
TiB2 is now considered as a promising reinforcements candidate for steels because of its high microhardness, excellent Young’s modules1) and good transverse rupture strength.2) In addition, TiB2 is relatively stable in liquid Fe3–5) and the solubility of Fe in TiB2 is lower than 4%.6) Besides, liquid Fe can well wet TiB2.7,8) More importantly, Fe–TiB2 composites fabricated by powder metallurgy (PM) and casting both show good interfacial cohesion between the Fe matrix and TiB2 reinforcements.9–13) Of most reported Fe–TiB2 composites, they are expected to apply to high modulus steels6,11,14–24) and wear-resistance parts.25–30) However, limited literatures reported about the application on hot work tools (HWTs).31)
In fields of HWTs, both good wear resistance and high thermal conductivity are necessary, since improving wear resistance can prolong the service time of HWTs and increasing thermal conductivity can bring an order of magnitude increase in thermal fatigue resistance32) and accelerate heat transfer from workpieces to cooling system. And thus, the higher production efficiency can be expected because the cycle time of forming is reduced. Moreover, TiB2 is endowed with high thermal conductivity of 96 W/(m·K),33) while it is also reported of 65–120 W/(m·K).1) Therefore, TiB2 reinforced Fe matrix composites are considered attractive materials for the usage not only in high modules steels and wear-resistance parts but also in HWTs due to the improved wear resistance and thermal conductivity.
In Fe–TiB2 system, Fe2B and TiC are commonly observed resultants resulted from the reaction occurring between Fe and TiB2. The reaction mechanism remains unclear. Impurities of C and O involved from starting powders are considered responsible for the decomposition of TiB2 in Fe, since impurities can consume Ti by formation of TiC or TiO2, which lead to the excess B. Then B atoms react with Fe, giving rise to the formation of Fe2B.34) However, in the absence of C and O, formation of Fe2B is still reported in literatures.28,35) The reaction mechanism of Fe and TiB2 is of great importance with regard to the synthesis of Fe–TiB2 composites, and hence it will be discussed in this study. In most cases, formation of Fe2B is undesirable because it may deteriorate toughness of materials, however, wear resistance of composites can be enhanced significantly.36,37) As for the formation of TiC, this is also, to some extent, beneficial to reinforce the matrix by TiC–TiB2 dual phase.38,39) As orientation relationship (OR) exists both in Fe–TiC ((101)Fe // $(1\bar{1}1)_{TiC}$)6,11) and TiC–TiB2 ({111}TiC // $(0001)_{TiB_{2}}$,40) which can contribute to a good interfacial cohesion between the interfaces.
Various synthesis routes have been put into place to fabricate Fe–TiB2 composites.12,19,26–29) Casting is a conventional and efficient route which can guarantee a clean interface, allow the tailoring of added elements as well as capable of casting large size ingots. However, there are two prominent issues which are clustering of TiB2 reinforcements in Fe matrix12) and density-induced floatation of the primary TiB2 particles.19) Apparently, more elaborate process should be taken into consideration in order to obtain a homogenously TiB2-distributed Fe matrix composites by casting route. Other processes involving in in-situ chemical reactions, most of them is self-propagating high temperature synthesis which enables a clean interface and allow adjustment of fraction of reinforcements. Nevertheless, the salient issues are difficult to control the process, aggregate of reinforcements and high porosity of resultants.26–29) The stereotyped disadvantages of PM process include the limitations to small parts and unclean interface between the matrix and the reinforcements. However, spark sintering technique, one novel method of PM, has been developed to be a high-efficiency sintering technique in recent decades. It can be applied to synthesizing a wide range of materials, such as metals,41) alloys,42) ceramics43) and composites.44) Due to the discharges, the particle surface is activated and purified,45–47) and a self-heating phenomenon is generated between the particles. As a result, heat-transfer and mass-transfer can be completed instantaneously. Therefore, spark sintering technique can be applied to sinter metal matrix composites quickly to the full density at a relatively low temperature. In addition, a homogenous microstructure can be obtained as long as a suitable mixing process is applied.
The objective of this study is to investigate the microstructures of Fe–TiB2 composites sintered by spark sintering and to fabricate the Fe–TiB2 composites with both high thermal conductivity and hardness for the possibility of HWTs applications.
As-received TiB2 (99.9%, Pure Chemical Co., Ltd. Japan) and pure Fe (99.9%, Sanwa Metal Industry Co., Ltd. Japan) powders with size of 2∼3 µm and 3∼5 µm, respectively, were used as starting powders. The wet mixing process was utilized to mix the 30 vol%TiB2 and 70 vol%Fe powders. TiB2 and Fe powders of total amount of 15 g were placed in a stainless-steel jar with a volume of 250 cm3 and then 25 mL ethanol was added as a wet mixing agent. Hereafter, the jars were vacuumed for 0.3 ks. Stainless-steel balls with a diameter of 2∼10 mm were used for mixing. The weight ratio of balls to powders was 10:1. Mixing process was conducted in a planetary ball mill (Pulverisette-5, Fritsch, Germany) at 100 rpm for 10.8 ks. Slurry of powder mixtures were dried in a fume hood and then dry mixed for 3.6 ks. Powder mixtures were loaded in a graphite die and consolidated by spark sintering (CS12567, NSK. Ltd. Chūgoku Branch, Japan) technique. All spark sintering experiments were carried out at an applied pressure of 50 MPa and a vacuum condition (<10−2 Pa). The die temperature was monitored using R type thermocouple inserted into the hole of die. Sintering temperature was determined at 1373 K because it was the optimized temperature based on previous work. Therefore, samples were sintered at 1373 K, heating rate of 100 K/min, and holding different times (0, 0.3, 0.6, 1.8 and 3.6 ks). Finally, compacts with height of 10 mm and diameter of 10 mm were obtained.
2.2 Characterizations of starting powders and sintered compactsMorphology of powders and microstructure of sintered compacts were observed by electron probe micron analyzer (EPMA, JXA-8900, JEOL, Japan). Phase compositions of sintered compacts were characterized by X-ray diffraction (XRD, D/max-2500PC, Rigaku, Japan) with Cu Kα radiation (L = 0.15418 nm, 40 kV, 100 mA). Scanning transmission electron microscope (STEM, Talos F200X, Thermo Fisher Scientific, USA), high-resolution TEM (HRTEM) and energy disperse spectroscopy (EDS) were performed at 200 kV. High angle annular dark field (HAADF) images and bright field (BF) images were taken at angle of 59–200 mrad and 9 mrad, respectively. Vickers hardness was obtained by Vickers hardness tester (FV-110, Future-Tech Corp., Japan) at a load of 5 kg and a dwelling time of 10 s. The thermal diffusivities (α) of the samples (Φ10 × 1 mm) were measured at room temperature by a laser flash thermal constants measuring apparatus (Thermal conductivity-9000h, Ulvac-riko, Japan). The specific heat capacities (Cp) of the samples (Φ4 × 0.5 mm) were measured at room temperature by differential scanning calorimeter (STA449C, Netzsch-Gerätebau GmbH, Germany). Densities (ρ) of sintered compacts were acquired by Archimedes’ principle. Thermal conductivities (λ) of sintered compacts were calculated according to the eq. (1). Image analysis was adopted to roughly calculate the area fraction of each phase in sintered compacts. Different phases were dyed in different colors and then the area fraction of each phase could be calculated by image analysis.
| \begin{equation} \lambda = \alpha \cdot C_{\text{p}} \cdot \rho \end{equation} | (1) |
Morphologies of starting powers are shown in Fig. 1. As-received pure Fe particles were spherical with size of 3∼5 µm as shown in Fig. 1(a). As-received TiB2 particles basically presented as prismatic particles, such as a hexagonal shape as shown in Fig. 1(b). Large hexagonal TiB2 particles with size of about 2∼3 µm corresponded to well-grown TiB2. In contrast, some kinds of small TiB2 were prismatic particles with size of approximate 500 nm. Powder mixtures of Fe–30 vol% TiB2 after mixing was shown in Fig. 1(c), and the magnified image of the area marked by white square in (c) was shown in Fig. 1(d) in which TiB2 and Fe were indicated according to the above described characteristics.
SEM images showing as received (a) Fe and (b) TiB2 powders, (c) powder mixtures of Fe–30 vol% TiB2 after mixing, and (d) high magnified image corresponding to the area marked by white square in (c).
XRD patterns of Fe–30 vol% TiB2 compacts sintered at 1373 K with different holding times are revealed in Fig. 2. Except the original phases of Fe and TiB2, newly formed phases of Fe2B and TiC were also found in all sintered compacts due to the reaction between Fe and TiB2. The C source was introduced from the graphite dies. According to Fe–C binary phase diagram, the C content in γ-Fe is about 2.08 mass%.48) This C source resulted in the formation of TiC.
XRD patterns of Fe–30 vol%TiB2 compacts sintered at 1373 K for (a) 0, (b) 0.3, (c) 0.6, (d) 1.8, and (e) 3.6 ks, respectively.
The typical back scatter electron (BSE) images of compacts sintered for 0, 0.6 and 3.6 ks are shown in Fig. 3. According to point analysis in our previous work,49) black areas, gray areas, light areas and circle-like areas corresponded to TiB2, Fe2B, Fe (or Fe solid solution with about 3 at%Ti) and TiC, respectively. This result coincided well with that from XRD patterns (shown in Fig. 2). TiB2 particles were basically well-dispersed in matrix. Fe2B appeared to be irregular morphology and was distributed throughout the whole compacts. TiC particles displayed surrounding TiB2 particle. Nevertheless, the areas between TiC circles and TiB2 was still yet to know but will be revealed in the subsequent STEM images. The area fraction of Fe2B increased with extending holding time, in contrast, that of Fe decreased because of the consumption of Fe by B atoms, as it could be appreciably distinguished on BSE images. In order to quantitatively obtain the fraction of each phase, the area fraction of each phase in all sintered compacts was roughly measured by image analysis as shown in Fig. 4(a). With the increase of holding time, the area fractions of TiB2 (24.4%–19.0%) and Fe (39.4%–30.0%) were decreasing while that of Fe2B (29.6%–42.0%) and TiC (6.6%–9.0%) were increasing as the reaction deepened with the increase of holding time. As shown in Fig. 4(b), the relative density was significantly improved from 0–0.6 ks, while it varied slightly from 0.6–3.6 ks. The improvement of relative density was mainly due to the elimination of sintering pores with the increase of holding time. The variation trend of relative density could be interpreted that holding time of 0.6 ks was effective to remove the most pores during sintering.
Typical BSE images of Fe–30 vol%TiB2 compacts sintered at 1373 K for different times: (a) 0, (b) 0.6, and (c) 3.6 ks, respectively.
Area fractions of each phase in (a) and relative density in (b) of Fe–30 vol%TiB2 compacts sintered at 1373 K for 0, 0.3, 0.6, 1.8 and 3.6 ks, respectively.
HAADF images and corresponding mapping analysis images of each element provided a better observation of phases distribution. Figure 5 corresponds to the compact sintered at 1373 K holding 0 ks. Figure 5(a)(b) are HAADF images and Fig. 5(b) is the high magnification image corresponding to the marked area in Fig. 5(a). Figure 5(c)(d)(e)(f) placed on the right are mapping analysis images of Fe, Ti, B and C, which corresponds to Fig. 5(a). Coupled with mapping analysis, TiB2, Fe2B, Fe and TiC were indicated in Fig. 5(a)(b). It was noticeable that the areas, mentioned in Fig. 3, connecting TiB2 and TiC particles were proved to be Fe as marked in Fig. 5(b). Namely, TiB2 was wrapped by a layer of Fe after sintering. This phenomenon will be interpreted in the following section of reaction mechanism between TiB2 and Fe. The preferred interface planes for TiB2 particles are prismatic $\{ 10\bar{1}0\} $ planes and the basal {0001} planes in Fe–TiB2 composites fabricated by casting.11) However, in present study, interfaces between Fe and TiB2 (Fe/TiB2 interface) were mainly parallel to the prismatic planes of TiB2 as marked in Fig. 5(b). As shown in Fig. 5(b), Fe impregnated the triple junctions between Fe and TiB2 as marked by white circles, and the interfaces of Fe and TiB2 were well bonded as marked by dashed arrows, which both demonstrated a good cohesion of Fe/TiB2 interface. Moreover, the good cohesion of Fe/TiB2 interface will be further characterized at atomic level in Fig. 6. TiC grains precipitating along Fe grain boundaries were arrayed in circles as evidenced in Fig. 5(a) and it could be further confirmed by the mapping analysis as marked in Fig. 5(e)(f).
HAADF images of Fe–30 vol%TiB2 compacts sintered at (a) 1373 K for 0 ks and (b) the high magnification image corresponding to the area in (a) marked by dashed box; (c), (d), (e) and (f) placed on the right are mapping analysis images of Fe, Ti, B and C corresponding to (a).
(a) A HRTEM image of Fe/TiB2 interface in the compact sintered at 1373 K for 0 ks, and an inset in the lower right corner is an FFT image corresponding to the area marked in dashed box, and (b) an IFFT image of Fe/TiB2 interface corresponding to the area marked in dashed box in (a).
In the light of work conducted by Cha6) and Korine,11) fantastic interpretations of Fe/TiB2 interface at an atomic level obtained by eutectic solidification are evidenced, providing a perspective of better understanding to the interfacial cohesion of Fe/TiB2. However, little literatures report the Fe/TiB2 interface of Fe–TiB2 composites obtained by PM. Therefore, High-resolution characterizations at microscopic and atomic levels were conducted to characterize the Fe/TiB2 interface in Fe–TiB2 composites synthesized by PM. A HRTEM image taken on Fe/TiB2 interface and a fast Fourier transform (FFT) image located in lower right corner corresponding to the marked area are shown in Fig. 6(a). In order to evaluate the Fe/TiB2 interface, the FFT image was filtered and inverse fast Fourier transformed (IFFT) as exhibited in Fig. 6(b). Atomic columns were visible in both sides. (100) plane of Fe was parallel to $(10\bar{1}0)$ prismatic plane of TiB2 and $(\bar{1}01)$ plane of Fe was inclined by 7° from $(01\bar{1}1)$ of TiB2. The misfit in the interface between $(01\bar{1}1)$ plane of TiB2 and $(\bar{1}01)$ plane of Fe was 5.9%. It should be accommodated by periodically spaced dislocations which could not be clearly evidenced due to contrast variations in interface, caused by defocus and thickness. This OR guaranteed a good interfacial cohesion between two phases. As it could be assumed that the prismatic planes of TiB2 are probably served as nucleation sites for α-Fe transformation from γ-Fe during cooling in term of the occurrence of the special OR.6,11) A good interfacial cohesion of Fe/TiB2 interface synthesized by PM method was also confirmed at an atomic level.
Figure 7 shows the HAADF images of compacts sintered at 0.6 and 3.6 ks. TiB2, Fe2B, Fe and TiC were indicated, as shown in Fig. 7(a)(b). Fe filled in the gap with width of approximate 50 nm between two TiB2 particles according to the scale bar as shown in Fig. 7(a) and no sintering pores were observed at the Fe/TiB2 interfaces, which demonstrated an excellent interfacial cohesion between Fe and TiB2. The preferred interface planes for TiB2 were prismatic planes as indicated in Fig. 7(a), but they were not distinguishable in Fig. 7(b) as marked by round-shaped TiB2 due to the overreaction between Fe and TiB2 when sintered for 3.6 ks.
HAADF images of Fe–30 vol%TiB2 compacts sintered at 1373 K for (a) 0.6 and (b) 3.6 ks.
It is seemingly surprising to be able to observe dislocations in TiB2 particles as shown in Fig. 8 corresponding to the compacts sintered at 1373 K holding 0.6 ks. Not only in this work, other findings about the observation of dislocations in TiB2 particles were also published. J.R. Ramberg et al.50) studied the temperature deformation of polycrystalline TiB2 at a strain rate of 5 × 10−4 s−1 and revealed dislocation glide was the deformation mechanism. Wang et al.51) conducted compression test on NiAl–TiB2 composites at various strain rates at 1140 and 1300 K and observed the high density of dislocation in TiB2 particles. Besides, the yield stress at a strain rate of 1 × 10−4 s−1 was about 70 MPa for 10–20 vol%TiB2–NiAl composites. J.D. Whittenberger et al.52) also reported the dislocation observation within TiB2 particles from 30 vol%TiB2–NiAl composites synthesized by hot pressing at 1675 to 1775 K and 128 MPa. Lartigue-Korinek discovered that both basal and prismatic slip plane of TiB2 were activated concerning the specimen subject to hot rolling at 1173 K.11) Therefore, ceramics tend to produce dislocations at high temperature and stress. In the case of compacts sintered at 1373 K, 50 MPa, it is quite possible to be able to observe dislocations in TiB2 particles.
A BFTEM image of Fe–30 vol%TiB2 compact sintered at 1373 K for 0.6 ks.
It is quite common to observe dislocations in metals while this is not the case where occurs a lot in ceramics, especially at low temperature. Because ceramics are devoid of slip systems and predominated by directional bonding as opposed to that in metals, consequently, ceramics cannot accommodate deformations properly and is prone to be fragile at low temperature. Therefore, it is worth noticing the occurrence of deformation behavior in TiB2 particles during spark sintering. This occurrence of deformation maybe contribute to the activation of slip system of TiB2 particles at high temperature and stress. In the work of J.R. Ramberg,50) five independent slip systems are required to accommodate a large deformation in polycrystalline. Both basal and prismatic slip systems are activated since these are the most easily activated slip systems, which have also been confirmed by Lartigue-Korinek.11) The pyramidal slip system $\{ 10\bar{1}1\} $ $\langle 11\bar{2}0\rangle $ would not take into consideration since it is generally activated only at very high stresses. $\{ 10\bar{1}0\} $ ⟨0001⟩, {0001} $\langle 10\bar{1}0\rangle $ and $\{ 11\bar{2}0\} $ $\langle 1\bar{1}00\rangle $ are considered to be the remaining possible three slip systems. Moreover, the nature of the chemical bonding of TiB2, a mixture of covalent and metallic bonding characters, is also thought responsible to a plastic deformation of TiB2.11) Occurrence of plastic deformation behavior in TiB2 particles may be beneficial to the application of HWTs.
3.4 Reaction mechanism between TiB2 and FeThe formation of Fe2B and TiC is attributed to the reaction of TiB2 and Fe. But the mechanism of reaction is still under debate. Impurities of C and O in starting powders can consume Ti and lead to the excess B which react with Fe, giving rise to the formation of Fe2B.34) Whereas, in the authors’ previous work,49) Fe2B and TiC were not found in compacts sintered at 1273 K, although there was C existence. Meanwhile, both Fe2B and TiC were verified in compacts sintered at 1323 K and above. This contradictory, to the authors’ knowledge, derived from unspecified synthetic temperature. Dissolution of TiB2 in Fe is varying at different temperatures. Since negligible TiB2 dissolute in solid Fe when temperature is lower than 1273 K and hence, no Fe2B and TiC are generated. On the contrary, there is about 16 mass% solubility of TiB2 in liquid Fe at 1873 K.53) In addition, the Gibbs free energy of reactions were calculated using the thermodynamic data.54) TiB2 was more energetically favorable compared with Fe2B and TiC as shown in Fig. 9. But still, formation of Fe2B is reported in Fe–Ti–B system at high temperature55) even at the absence of C or O. This is ascribed to the heterogeneous distribution of Ti and B atoms in Fe, which will be interpreted later in terms of diffusion kinetics. Therefore, it is considered that dissolution of TiB2 in the Fe alloys lead to the formation of Fe2B.53) Another reason is negligible solubility of B in either α-Fe or in γ-Fe,55) resulting in a direct reaction between Fe and B after the decomposition of TiB2.
Gibbs free energy change of reactions as a function of temperature: (a) 2Fe + B → Fe2B, (b) Ti + C → TiC, and (c) Ti + 2B → TiB2.
As for the distribution of Fe2B and TiC in compacts sintered in 1373 K described aforehand in Fig. 3. This microstructure in Fig. 3 could be attributed to the different diffusion behavior of Ti and B driven by chemical potential. The diffusion mode of B atoms in Fe is mainly interstitial diffusion,56) but it is substitutional diffusion for Ti atoms in Fe.57) Some assumptions were made to approximately calculate the diffusion depth of Ti and B in Fe. Firstly, the diffusion model were simplified to be an infinite diffusion couple. According to Fick’s second law:
| \begin{equation} \frac{\partial \text{C}_{i}}{\partial t} = \text{D}_{i} \frac{\partial^{2}\text{C}_{i} (\text{x},\text{t})}{\partial x^{2}}. \end{equation} | (2) |
| \begin{equation} \text{D}_{i} = \text{D}_{0} \exp \frac{\text{$-$Q}}{\text{RT}}. \end{equation} | (3) |
| \begin{equation} \text{D}_{B}^{\textit{Fe}} = 4.4 \times 10^{-8} \exp \left( -\frac{81.5 \times 10^{3}}{\text{RT}} \right) \end{equation} | (4) |
| \begin{equation} \text{D}_{\textit{Ti}}^{\textit{Fe}(2\textit{mass}\%\textit{Ti})} = 0.56 \times \exp \left( -\frac{216.1 \times 10^{3}}{\text{RT}} \right). \end{equation} | (5) |
| \begin{equation} \text{x} \approx \sqrt{\text{Dt}}. \end{equation} | (6) |
According to eqs. (4),58) (5)59) and (6), approximate diffusion depth of B and Ti in Fe could be obtained if the compact was sintered at 1373 K for 0.6 ks, they were xB = 145 µm, and xTi = 14.4 µm, respectively. Apparently, the diffusion calculations employed were subject to error owing to over simplified diffusion model and to the inaccuracy of the diffusion data employed from literatures. However, it could be concluded that diffusion depth of B was an order of magnitude higher than that of Ti in Fe. The limited dissolution of TiB2 in Fe at 1273 K and below was due to the insufficient diffusion activation energy needed for atoms diffusion.
In the case of present work, the intrinsic modeling was radial diffusion for Ti and B in Fe. On the purpose of better interpreting the reaction mechanism between Fe and TiB2, schematic illustrations were drawn as shown in Fig. 10 based on the microstructure in Fig. 5. The whole process was roughly divided into four stages. In the first stage, as shown in Fig. 10(a), it showed TiB2 particle was surrounded by Fe particles before the reaction occurred between Fe and TiB2. The second stage was atom diffusing. According to our previous study,49) TiB2 particles started to decompose in pure Fe when sintering temperature was over 1323 K. As displayed in Fig. 10(b), Ti atoms mainly distributed around TiB2 particle due to limited diffusion depth, while B atoms scattered more extensively than Ti because of longer diffusion depth, as explained in the above paragraph. Moreover, the further away from TiB2 particles, the lower concentration of Ti and B atoms were. Therefore, higher concentration of atoms was represented by larger icons. At the same time, Fe also diffused toward TiB2, which led to a narrow interdiffusion region as indicated by two dashed circles shown Fig. 10(b)(c). B atoms directly reacted with Fe after diffusing through the interdiffusion region, resulting in the formation of Fe2B, as demonstrated in Fig. 10(c). This is due to the solubility of B is negligible in both α-Fe and γ-Fe.55) As time went on, Fe2B grains grew up as shown in Fig. 10(d). The interdiffusion region was considered as a solid solution of Fe–Ti–C (Fess), so Ti and C atoms probably started to precipitate at Fe grain boundaries during the γ-α transformation in the fourth stage as shown in Fig. 10(d). This is because interface planes between Fe and TiC such as $(1\bar{1}0)_{Fe}$ // $(0\bar{2}0)_{TiC}$ or (101)Fe // $(1\bar{1}1)_{TiC}$ can be envisaged as preferred orientations for the TiC precipitation,6) on the other hand, the solubility of C in α-Fe is much smaller than that in γ-Fe. TiC other than Fe3C precipitated from the Fess, as the affinity between Ti and C is higher than that between Fe and C. The residual Fe comprised of two types. As shown in Fig. 10(d), one wrapped around TiB2 after TiC precipitation from the Fess as indicated by Fe(1); The other type was unreacted Fe as marked by Fe(2) which was initially far away from TiB2. Figure 10(d) is the final stage which can be proved by the microstructures of compacts as shown in Fig. 5 and 6.
Schematic illustrations of reaction mechanism between Fe and TiB2: (a) the first stage, (b) second stage, (c) third stage, and (d) fourth stage; In (b) and (c), the triangles, circles, pentagons and squares represent C, Ti, B and Fe atoms in Fess, respectively, and the larger icons represent higher concentration of atoms.
Microstructure of materials can dramatically influence the mechanical and thermal properties such as strength, hardness, and thermal conductivity. Simultaneously, these properties can determine the applications of these materials in industrial practice. The thermal conductivity and Vicker hardness of all compacts as a function of holding time was plotted in Fig. 11. The thermal conductivity of pure Fe, TiB2 Fe2B and TiC are reported to be 80, 100,1) 3560) and 2161) W/(m·K), respectively. Therefore, the theoretical thermal conductivity of sintered compacts were estimated by simple mixture rule based on the data as shown in Fig. 4. The variation trend of theoretical and measured curves were basically same as shown in Fig. 11(a). In addition, theoretical thermal conductivity of sintered compacts was decreased as sintering time increased, and was overall higher than the measured one because the defects excluding pores and interface thermal resistance between matrix and reinforcements were not taken into consideration in such calculation. However, defects in compacts62) and interface thermal resistance between matrix and reinforcements63) are two important factors influencing thermal conductivity of materials. As defects and interfaces scatter electron or phonon during heat transfer, which decrease the mean free length of the path of electron or phonon and thus lead to the decrease of thermal conductivity. Besides, in this study, newly formed phases with low thermal conductivity were another main reason affecting the whole thermal conductivity of composites. The thermal conductivity increased first then decreased as the holding time extended, reaching the maximum at 0.6 ks. Two stages could be roughly divided to explain variation of thermal conductivity. One was the densification stage where defects were gradually diminished from 0–0.6 ks resulting in the improvement of thermal conductivity. It could also be proved by relative density variation in Fig. 4(b) showing that the relative density was greatly improved from 0–0.6 ks while it kept almost constant from 0.6–3.6 ks. The other stage was the formation stage of large amount of phases with low thermal conductivity, giving rise to the decrease of thermal conductivity from 0.6–3.6 ks. As a consequence, thermal conductivity of compact at 0.6 ks showed a maximal value, which was probably a trade-off between densification and growing amount of phases with low thermal conductivity. Besides, the thermal conductivity of each sintered compact was also higher than that of SKD61. However, with the increase of holding time, the Vickers hardness increased dramatically from 0–0.6 ks, after that, it slightly increased from 0.6–3.6 ks as displayed in Fig. 11(b). This was attributable to the growing amount of hard phases (Fe2B and TiC shown in Fig. 4(a)) and to the improved densification shown in Fig. 4(b). It was concluded that the compact sintered at 1373 K for 0.6 ks showed the most favorable properties among all the compacts on the basis of relative density, thermal conductivity and hardness.
Theoretical and measured thermal conductivity in (a) and Vickers hardness in (b) of Fe–30 vol%TiB2 compacts sintered at 1373 K for 0, 0.3, 0.6, 1.8 and 3.6 ks, respectively.
The formation of Fe2B and TiC can assuredly improve the hardness of the sintered compacts, especially Fe2B which has been widely used in surface boronizing to enhance the wear resistance of steels. In the case of HWTs, both wear resistance and toughness are indispensable even though they somehow contradict to each other. Therefore, well-tailored microstructure still need to be investigated in order to reconcile the two properties in the future. But so far, Fe–TiB2 composites with both high thermal conductivity and hardness have been synthesized. Especially, the compact sintered at 1373 K for 0.6 ks exhibited the most favorable properties which were 133% higher in thermal conductivity and comparable in Vickers hardness compared with that of SKD61. Therefore, this work provides a new way to fabricate a future generation of HWTs as shown in Fig. 12.
Comparison between Fe–30 vol%TiB2 sintered compacts in present study and SKD61 in Vickers hardness and thermal conductivity.
The phase composition was Fe, TiB2, Fe2B and TiC in all sintered compacts. With the increase of holding time, the area fractions of TiB2 and Fe were decreasing while that of Fe2B and TiC were increasing. A good Fe/TiB2 interfacial cohesion was confirmed at atomic level, which was the occurrence of the special OR between {110} planes of Fe and $\{ 10\bar{1}0\} $ planes of TiB2. Dislocation observation in TiB2 particles proved the plastic deformation ability of TiB2 at high temperature and stress. Reaction between Fe and TiB2 was due to the decomposition of TiB2 in Fe at 1373 K and different diffusion depth of B and Ti in Fe. Thereafter, B directly reacted with Fe, since the solubility of B atoms is low in either α-Fe or in γ-Fe. TiC precipitated from Fess along Fe grains boundary as temperature dropped, which led to the residual Fe wrapping around TiB2. With the increase of holding time, the thermal conductivity increased first then decreased reaching the maximum at the holding time of 0.6 ks, which caused by a trade-off between increased densification and growing amount of phases with low thermal conductivity; However, Vickers hardness was increasing which was attributable to the growing amount of hard phases (Fe2B and TiC). Among all the compacts, the one sintered at 1373 K for 0.6 ks showed the excellent properties which were 133% higher in thermal conductivity and comparable in Vickers hardness compared with that of SKD61. This work provides a new way to fabricate a future generation of HWTs.
We gratefully acknowledge the support from LETS Research Group, Incubation Research Center of Hiroshima University, Y-TEC Corporation, KEYLEX Corporation, MAZDA and HATACHI METALS, Ltd. And Dr. Wentao Hu in Yanshan University is gratefully acknowledged for the help on TEM characterization.
