Porous high entropy (Zr_0.2Hf_0.2Ti_0.2Nb_0.2Ta_0.2)B₂: A novel strategy towards making ultrahigh temperature ceramics thermal insulating

1. Introduction

Ultrahigh temperature ceramics are considered as candidate materials for extreme-environmental applications due to their excellent thermal stability, high strength and modulus, and good oxidation resistance [[1], [2], [3], [4], [5]]. In ultrahigh temperature thermal managements, besides materials that can efficiently dissipating heat, thermal insulating materials in the form of coatings and porous bulk also play important roles in preventing thermal damages. However, the high thermal conductivity of transition metal diborides based UHTCs makes their application as thermal insulating materials difficult [[6], [7], [8], [9], [10]]. The high thermal conductivity of transition metal diborides comes from the contributions of both electrons and phonons. For example, ZrB₂ has a total thermal conductivity as high as 100-140 W m^-1 K^-1, with 88-91 W m^-1 K^-1 from the phonon contribution and the rest from electrical contribution [11]. Thus both electrical and phonon contributions must be simultaneously controlled in order to reduce the thermal conductivity of transition metal diborides. Recent successes in high entropy (HE) materials have revealed that both electrons and phonons are scattered in this new class of materials such that their thermal conductivities are significantly reduced, which opens a new window to design novel insulating materials [[12], [13], [14]]. In light of the high entropy effects, a highly porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is designed as a novel ultrahigh temperature thermal insulating material. The basis for designing of this new insulating material is the expression of Kingery’s [15] model:

κ_p=κs(1-P_o) (1)

where κ_p is the thermal conductivity of porous materials, κ_s is the thermal conductivity of backbone material, P_o is the cross-sectional pore fraction, from which one can see that decreasing the thermal conductivity of backbone material and increasing the porosity are key factors in realizing thermal insulating. In addition to lattice distortion effect in HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ backbone material, controlling grain size to nanometer ranges also favors to reducing thermal conductivity. Our previous work has demonstrated the possibility of achieving low thermal conductivity ($\widetilde{0}$.39 W m^-1 K^-1) in high porosity (over 80%) HE (Zr_0.2Hf_0.2Ti_0.2Nb_0.2Ta_0.2)C [14]. For transition metal diborides based UHTCs, low thermal conductivity has not been reached due to the intrinsic higher thermal conductivities than transition metal carbides and the difficulty in preparing porous bulk materials. Thus, high porosity is needed to make porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ thermal insulating.

In peruse of literature, one will find that previous works on high entropy transition metal diborides are mainly focused on synthesis of powders and SPS sintering of bulk. For example, Gild et al. [16] synthesized a series of HE metal diboride ceramics by high energy ball milling and spark plasma sintering (SPS). These HE borides possess better hardness, high-temperature stability and oxidation resistance. Tallarita et al. [17] fabricated HE (Hf_0.2Mo_0.2Ta_0.2Nb_0.2Ti_0.2)B₂ bulk ceramic by a two-step method consisting of self-propagating high-temperature synthesis (SHS) and SPS. Zhang et al. [18] prepared three types of HE metal diboride powders and dense ceramics through borothermal reduction and SPS. In addition, HE (ZrVTiTaHf)B₂ coatings were also prepared by Mayrhofer et al. [19] using non-reactive magnetron sputtering. Obviously, these methods are not suitable for preparing highly porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ with low thermal conductivity.

In order to prepare porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂, we have designed an in-situ borocarbon thermal reduction/partial sintering process as described in Eq. (2).

14ZrO₂ (s) + 14HfO₂ (s) + 14TiO₂ (s) + 7Nb₂O₅ (s) + 7Ta₂O₅ (s) + 52B₄C (s) = 70(Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ (s) + 34B₂O₃ (g) + 52CO (g) (2)

In this process, HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ can be synthesized at relatively low temperatures compared to the previously used methods, while partial sintering can be completed by subsequent heating at higher temperatures. High porosity can be formed by reaction produced B₂O₃ and CO gases. As we have demonstrated in our previous works [[20], [21], [22]], the advantages of this method includes: (1) low cost starting materials, (2) low synthesizing temperature, (3) no pore forming agent and no sintering additive are needed, (4) partial sintering is completed immediately after synthesis. Furthermore, selecting boron carbide (B₄C) as the reductive agent in this work instead of boron (B) is beneficial to increase the porosity of porous UHTCs, because borocarbon reduction (2) can release more gases than boron reduction (3).

6ZrO₂ (s) + 6HfO₂ (s) + 6TiO₂ (s) +3Nb₂O₅ (s) + 3Ta₂O₅ (s) + 104B (s) = 30(Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ (s) + 22B₂O₃ (g) (3)

In the following part of this letter, the shrinkage during reaction synthesis and sintering is investigated first. Then high porosity HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is prepared. The phase composition and microstructure of HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ are investigated. Finally, low thermal conductivity and high compressive strength are demonstrated.

2. Experimental

2.1. Synthesis of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂

The raw materials used in the present work were ZrO₂, HfO₂, TiO₂, Nb₂O₅, Ta₂O₅ (99.9%, 1 μm) and B₄C (99%, 2 μm) powders. According to Eq. (2), reactant powders were mixed to prepare the ethanol-based slurry. The slurry was ball-milled in a polyurethane jar with zirconia grinding balls for 10 h. After ball-milling, the slurry was dried in an oven at 60 °C until the ethanol evaporated completely. Then, the mixed powders and zirconia balls were separated by using a 120-mesh screen.

In order to determine the synthesis temperature of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ based on reaction (2), the reaction process was monitored using a thermal analyzer (NETZSCH STA 409 CD/7/403/5/G, Germany) through heating reactant powders to 1650 °C in flowing argon atmosphere at a heating rate of 5 °C /min. Meanwhile, shrinkage behavior during the heat-treatment was monitored using a thermal dilatometer (NETZSCH DIL 402 E, Germany) through heating the columnar shaped green bodies (7 mm in diameter and 15 mm in height) up to 1850 °C in flowing argon atmosphere at a heating rate of 5 °C/min.

For high porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ preparation, the reactant powders were uniaxially pressed into columnar shaped green bodies of 12 mm in diameter. Then the green body was heated to 1700 °C for HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ synthesis and to 1850 °C for 1 h in vacuum for partial sintering.

2.2. Characterization of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂

The density of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ was determined geometrically, wherein the theoretical density of (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ was calculated from the mass of the atoms in a unit cell and the lattice parameters measured from the XRD pattern of HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂. The porosity (P) of porous bulk was calculated according to Eq. (4):

P=1- ρρ'×100% (4)

where ρ is the density of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂, ρ′ is the theoretical density of HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂.

The formation of (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ solid solution was identified by X-ray diffraction analysis (XRD, D8 advanced, Bruker, Karlsruhe, Germany) using CuKα radiation (1.54178 Å). The pore size, grain size and element distribution of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ were analyzed in a scanning electron microscope (SEM, Apollo 300, CamScan, Cambridge, UK) with an energy dispersive spectroscopy (EDS) system. The compressive strength was measured by using a universal testing machine (DZS-III, China Building Material Test and Certification Center, Beijing, China) using rectangular bars of 5 mm × 5 mm × 10 mm in size.

Thermal insulation of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is characterized by its low thermal conductivity (κ), which was calculated according to Eq. (5):

κ=α·ρ·C_p (5)

where α is the thermal diffusivity, ρ is the density, C_p is the specific heat capacity. The parameter α was measured using a laser thermal conductivity testing instrument (ANTER FLASHLINE 5000, USA) with a sample size of Φ10 mm × 2 mm, while C_p was measured on a differential scanning calorimeter (DSC 250, TA Instruments, USA).

3. Results and discussion

3.1. Shrinkage history during heating

Since a large amount of B₂O₃ and CO gases will be produced in reaction (2), the shrinkage must be carefully controlled to avoid cracking during reaction synthesis. Thus, the shrinkage during heating process must be monitored as we have done in our previous works [14,21]. Fig. 1(a) shows the linear shrinkage (dL/L₀, with L₀ being the height of green body at room temperature) and linear shrinkage rate ((1/L₀)dL/dt, with L being the height of the sample, t being the time)) curves. In addition, since the borocarbon reduction reaction (2) is accompanied by mass loss, the change of mass must also be monitored to investigate the reaction process. Fig. 1(b) shows the thermal gravity (TG) and differential thermal gravity (DTG) curves. It can be seen that the reaction process for HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ formation can be recognized as starting at 870 °C and finishing at 1700 °C. The shrinkage at temperatures above 1700 °C is mainly due to sintering of HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂. Based on the shrinkage and mass loss history shown in Fig. 1, the preparation parameters for porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is selected as follows: (1) rapid rate heating ($\widetilde{2}$0 °C/min) to 870 °C; (2) slow rate heating (under 5 °C/min) between 870-1700 °C; (3) partial sintering at 1850 °C. Herein, slow rate heating between 870-1700 °C is selected to avoid cracking caused by rapid shrinkage and gas-releasing, while relatively low sintering temperature of 1850 °C is selected to avoid grain growth and ensure high porosity.

3.2. Microstructure and compressive strength

Using the preparation parameters determined in the above section, porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ without cracks has been prepared. As listed in Table 1, the density of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is 2.01 g/cm³. Fig. 2(a) shows the X-ray diffraction (XRD) pattern of as-synthesized HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ together with simulated XRD patterns of ZrB₂, HfB₂, TiB₂, TaB₂ and NbB₂. The simulated XRD patterns of ZrB₂, HfB₂, TiB₂, TaB₂ and NbB₂ are generated from the Reflex code in Materials Studio software suit (Accelrys Software Inc., San Diego, CA, USA) using the crystal structure data of Ref. [23] and CuKα radiation of wavelength = 1.5418 Å, step size of 0.02° and Pseudo-Voigt peak shape function. From the XRD pattern, it can be seen that a single-phase solid solution (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ with the hexagonal AlB₂-type crystal structure is formed. The lattice parameters of the (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ solid solution are a = b =3.1036(1) Å, c = 3.3795(3) Å. Using the lattice constants and atomic mass as input, the theoretical density is calculated to be 8.26 g/cm³. Then the porosity is determined as 75.67%. High porosity of HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is attributed to the large amount of CO and B₂O₃ gases produced during in-situ reaction/partial sintering process, which prevent the densification and consequently develop a great number of pores in microstructure [[20], [21], [22]].

Fig. 2(b) shows the surface SEM image of as-prepared porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂. It can be seen that the as-prepared porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ possesses homogeneous microstructure with small grain size (400-800 nm), small pore size (0.3-1.2 μm) and high porosity. The porous structure is formed by necking between grains during partial sintering process. The corresponding EDS mappings of Zr, Hf, Nb, Ta and Ti are shown in Fig. 3. The distribution of these transition metal elements is homogeneous at the micron scale, which indicates that a solid solution is formed.

Based on the results of XRD and SEM-EDS analysis, the as-prepared porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is homogeneous in composition and structure. In addition, reaction liberated gas is thought to promote the neck growth between grains in the initial stage of sintering process [24,25]. As a result of necks between grains, the strength of porous ceramics can be enhanced [26,27]. As listed in Table 1, the porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ prepared by an in-situ borocarbon thermal reduction /partial sintering process exhibits high compressive strength of 3.93 MPa.

3.3. Thermal conductivity

In order to demonstrate that the effects of high entropy and porosity on the thermal insulating behavior, the thermal conductivity of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ is investigated. The thermal diffusivity, heat capacity and thermal conductivity of porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ are listed in Table 2. As expected, the as-prepared porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ exhibits low thermal conductivity in the temperature range of 298-500 K. The RT thermal conductivity of porous HE (Zr_0.2Hf_0.2Ti_0.2 Nb_0.2Ta_0.2) B₂ is 0.51 W m^-1 K^-1, which is only 1/100-1/50 of the porous ZrB₂ fabricated by Yuan et al. [28]. Fig. 4 summarizes the thermal conductivity versus porosity for dense and porous UHTCs [10,11,14,21,[28], [29], [30]]. It is obvious that HE (Zr_0.2Hf_0.2Ti_0.2 Nb_0.2Ta_0.2) B₂ has low thermal conductivity. It also demonstrates that designing porous HE UHTCs is a novel strategy in changing UHTCs from thermal conducting to thermal insulating.

4. Conclusion

A new thermal insulting ultrahigh temperature material, porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂, is designed and prepared by an in-situ thermal borocarbon reduction/partial sintering process in this work. A single-phase solid solution (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ with the hexagonal AlB₂-type crystal structure is formed. The as-prepared porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ possesses high porosity of 75.67%, small pore size of 0.3-1.2 μm, homogeneous microstructure with small grain size of 400-800 nm, high compressive strength of 3.93 MPa, and low RT thermal diffusivity and thermal conductivity of 0.74 mm² s^-1 and 0.51 W m^-1 K^-1, respectively. The combination of these properties makes porous HE (Zr_0.2Hf_0.2Nb_0.2Ta_0.2Ti_0.2)B₂ promising for ultra-high temperature thermal insulation applications.

Acknowledgment

This work was supported by the National Natural Science Foundation of China (Nos. 51672064 and U1435206).

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