Effects of HfB2 and HfN Additions on the Microstructures and Mechanical Properties of TiB2-Based Ceramic Tool Materials

The effects of HfB2 and HfN additions on the microstructures and mechanical properties of TiB2-based ceramic tool materials were investigated. The results showed that the HfB2 additive not only can inhibit the TiB2 grain growth but can also change the morphology of some TiB2 grains from bigger polygons to smaller polygons or longer ovals that are advantageous for forming a relatively fine microstructure, and that the HfN additive had a tendency toward agglomeration. The improvement of flexural strength and Vickers hardness of the TiB2-HfB2 ceramics was due to the relatively fine microstructure; the decrease of fracture toughness was ascribed to the formation of a weaker grain boundary strength due to the brittle rim phase and the poor wettability between HfB2 and Ni. The decrease of the flexural strength and Vickers hardness of the TiB2-HfN ceramics was due to the increase of defects such as TiB2 coarse grains and HfN agglomeration; the enhancement of fracture toughness was mainly attributed to the decrease of the pore number and the increase of the rim phase and TiB2 coarse grains. The toughening mechanisms of TiB2-HfB2 ceramics mainly included crack bridging and transgranular fracture, while the toughening mechanisms of TiB2-HfN ceramics mainly included crack deflection, crack bridging, transgranular fracture, and the core-rim structure.


Introduction
In recent years, with the widespread use of difficult-to-machine materials in engineering, cutting tools have faced the challenge of machining these materials under high speed, which requires that the tools have high hardness, excellent wear resistance, oxidation resistance, and so on. However, compared with the ceramic tool materials, the traditional tool materials (high-speed steel and cemented carbide) showed a lower red hardness in machining these difficult-to-machine materials, which did not meet the need of high speed machining. Recently, ceramic tools-Al 2 O 3 -based, Si 3 N 4 -based, and TiB 2 -based ceramic tools-exhibited excellent cutting performance in machining the difficult-to-machine materials such as martensitic stainless steel, Inconel 718, ultra-high-strength steel 300 M, heat-treated AISI4140, hardened Cr12MoV mold steel, and Invar36 alloy [1][2][3][4][5][6]. The TiB 2 -based ceramic tool exhibited higher hardness compared with the other ceramic tools, which was attributed to the higher hardness of TiB 2 ceramic than that of the other ceramics. TiB 2 also has a high melting point, excellent wear resistance, and oxidation resistance, which can also be applied in other fields such as manufacturing armor plates and dies [7][8][9][10]. However, it shows a tendency toward low flexural strength and low fracture toughness, which limits the more widespread application of TiB 2 . In order to reverse this tendency and improve the mechanical properties of TiB 2 ceramic, reinforcements such as hard phases, metal phases, and whiskers have been employed to fabricate TiB 2 -based ceramic materials through spark plasma sintering, vacuum hot-pressed sintering, or reactive hot-pressed sintering.
Usually, the hard phases included TaC, TiSi 2 , Al 2 O 3 , WC, TiC, B 4 C, NbC, MoSi 2 , SiC, and ZrB 2 [9][10][11][12][13][14][15][16], which could inhibit the grain growth of the base material to obtain a fine microstructure. HfB 2 and HfN have high hardness, high melting point, and high oxidation resistance, and as reinforcements they can enhance the mechanical properties of ceramics such as ZrB 2 -CrSi 2 -HfB 2 , ZrB 2 -SiC-HfB 2 , B 4 C-HfB 2 , and SiBCN-HfN [17][18][19][20], which make them potential candidate reinforcements for ceramic tool materials. In addition, because HfB 2 and HfN have better thermal stability to resist deformation and decomposition at elevated temperature, they may improve the cutting performance and working life of TiB 2 -based ceramic tools. The metal phases often contain Fe, Co, Ni, and Mo [7,12,21,22], which could decrease the sintering temperature and improve the boundary strength among grains and relative density, while the ceramic whiskers such as aluminum borate whiskers and SiC whiskers could change the direction of crack growth to consume more crack propagation energy [23][24][25], which could improve the flexural strength and fracture toughness. Usually, adopting a combination of reinforcements for fabricating TiB 2 -based ceramics can obtain better mechanical properties. In addition, compared with spark plasma sintering that is employed in fabricating the ceramic composites [9,16], vacuum hot-pressed sintering is considered to be easily adaptable and economically viable.
In this paper, TiB 2 -HfB 2 and TiB 2 -HfN ceramic tool materials will be fabricated with powders of TiB 2 , HfB 2 , HfN, Mo, and Ni by vacuum hot-pressed sintering. The characteristics of these composites are analyzed according to their microstructures and mechanical properties.

Experimental Procedures
Commercially available TiB 2 powder (99.9%, 1 µm, Shanghai Xiangtian Nanomaterials Co., Ltd., Shanghai, China), HfB 2 powder (99.9%, 0.8 µm, Shanghai Chaowei Nanomaterials Co., Ltd., Shanghai, China) and HfN powder (99.9%, 0.8 µm, Shanghai Chaowei Nanomaterials Co., Ltd.) were used as the raw materials. Ni powder (99.8%, 1 µm, Shanghai Yunfu Nanotechnology Co., Ltd., Shanghai, China) and Mo powder (99.8%, 1 µm, Shanghai Yunfu Nanotechnology Co., Ltd.) were added as sintering aids. The compositions of the composite tool materials are shown in Table 1. The powders were mixed and milled for 48 h in a polyethylene jar with WC (tungsten carbide) balls and alcohol as the medium. Then the mixed slurry was dried in vacuum and sieved by a 200-mesh sieve. The compacted powders were hot pressed for 30 min at 1650 • C under 30 MPa in a vacuum ((1.2-2.4) × 10 −3 Pa). The hot pressed samples were cut into testing specimens by the electrical discharge wire cutting method and the surfaces of the testing bars were polished using diamond slurries. The dimensions of the specimens were 3 mm × 4 mm × 40 mm.
Flexural strength was measured at a span of 30 mm and a crosshead speed of 0.5 mm/min by the three-point bending test method on an electron universal tester (CREE-8003G, Dongguan City Kerry Instrument Technology Co., Ltd., Dongguan, China), according to Chinese National Standards GB/T 6569-2006/ISO 14704:2000 [26]. The fracture toughness (K IC ) was measured via the direct indentation method and was calculated through the following equation [12,27]: where H V is the Vickers hardness, 2a is the length of the impression diagonal, and 2c is the overall indentation crack length including 2a.

Results and Discussions
3.1. Microstructure Figure 1 shows the XRD patterns of the TiB 2 -HfB 2 and TiB 2 -HfN ceramic tool materials. The major crystal phases are TiB 2 and HfB 2 in the TiB 2 -HfB 2 ceramics, and TiB 2 and HfN in the TiB 2 -HfN ceramics. The minor phase is the Ni 3 Mo intermetallic compound in the TiB 2 -HfB 2 and TiB 2 -HfN ceramic tool materials. This is because Ni and Mo can form the Ni 3 Mo intermetallic compound at 1300 • C [29]. The Ni 3 Mo intermetallic compound has a high melting point of about 1320 • C, so it may be a promising high-temperature structural material [30]. Compared with the standard peaks, the peaks of HfB 2 are offset about two degrees to the right and are near the peaks of TiB 2 . This indicates there is likely an exchange of Ti and Hf atoms in the sintering, which leads to a complex solid solution of TiB 2 and HfB 2 formed in the ceramic tool materials. The peaks of HfN are in accordance with the standard peaks.  Figure 2a-c looks like the TiB2 grain shape and the irregular pore shape in Figure 2d-f looks like the shape of agglomerated HfN grains possibly pulled out in the grinding and polishing process, which indicates that a weaker grain boundary strength formed in these ceramics in the sintering processing. Moreover, in Figure 2a-c the morphology of some TiB2 grains changes from bigger polygons to smaller polygons or longer ovals which is advantageous for the formation of a relatively fine microstructure, and in Figure 2d-f the HfN grain agglomeration becomes more serious leading to the formation of more TiB2 coarse grains and pores. This indicates that the HfB2 additive not only can inhibit the growth of TiB2 grains but can also change the morphology of some TiB2 grains, and that the HfN additive exhibits a tendency toward agglomeration.  Figure 2a-c looks like the TiB 2 grain shape and the irregular pore shape in Figure 2d-f looks like the shape of agglomerated HfN grains possibly pulled out in the grinding and polishing process, which indicates that a weaker grain boundary strength formed in these ceramics in the sintering processing. Moreover, in Figure 2a-c the morphology of some TiB 2 grains changes from bigger polygons to smaller polygons or longer ovals which is advantageous for the formation of a relatively fine microstructure, and in Figure 2d-f the HfN grain agglomeration becomes more serious leading to the formation of more TiB 2 coarse grains and pores. This indicates that the HfB 2 additive not only can inhibit the growth of TiB 2 grains but can also change the morphology of some TiB 2 grains, and that the HfN additive exhibits a tendency toward agglomeration.   Figure 4 shows the fracture morphology of the TiB2-HfB2 and TiB2-HfN ceramic tool materials. As can be seen in Figure 4a-c, with increasing HfB2 content from 10 wt % to 30 wt %, the TiB2 grains become smaller; meanwhile, the TiB2 grain shapes exhibit the same variation trend as presented in Figure 3a-c; moreover, the pore number decreases progressively. However, in Figure 4d-f with increasing HfN content from 10 wt % to 30 wt %, the TiB2 grains become larger leading to the formation of coarse TiB2 grains; and the pore number decreases progressively. The results indicate   Figure 4 shows the fracture morphology of the TiB2-HfB2 and TiB2-HfN ceramic tool materials. As can be seen in Figure 4a-c, with increasing HfB2 content from 10 wt % to 30 wt %, the TiB2 grains become smaller; meanwhile, the TiB2 grain shapes exhibit the same variation trend as presented in Figure 3a-c; moreover, the pore number decreases progressively. However, in Figure 4d-f with increasing HfN content from 10 wt % to 30 wt %, the TiB2 grains become larger leading to the formation of coarse TiB2 grains; and the pore number decreases progressively. The results indicate  Figure 4 shows the fracture morphology of the TiB 2 -HfB 2 and TiB 2 -HfN ceramic tool materials. As can be seen in Figure 4a-c, with increasing HfB 2 content from 10 wt % to 30 wt %, the TiB 2 grains become smaller; meanwhile, the TiB 2 grain shapes exhibit the same variation trend as presented in Figure 3a-c; moreover, the pore number decreases progressively. However, in Figure 4d-f with increasing HfN content from 10 wt % to 30 wt %, the TiB 2 grains become larger leading to the formation of coarse TiB 2 grains; and the pore number decreases progressively. The results indicate that the HfB 2 additive can not only inhibit the growth of the TiB 2 grains, but can also change the microstructure of TiB 2 -based ceramic, and that the HfN additive cannot inhibit the TiB 2 grain growth. that the HfB2 additive can not only inhibit the growth of the TiB2 grains, but can also change the microstructure of TiB2-based ceramic, and that the HfN additive cannot inhibit the TiB2 grain growth.  Figure 5 presents the relative densities of the TiB2-HfB2 and TiB2-HfN ceramic tool materials. As can be seen, their relative densities increase with increasing HfB2 and HfN contents from 10 wt % to 30 wt %, respectively. The relative density increments of the TiB2-HfN ceramics are smaller, and the relative density variation curve is relatively flat, while the relative density variation curve of TiB2-HfB2 shows a bigger increment in relative density at first, and then finally shows a smaller increment. These results are ascribed to the pore number reduction with increasing the additive content, to some extent, and is derived from the higher sintering pressure (30 MPa) and the metal phases (Ni and Mo) that can efficiently reduce the sintering temperature and can accelerate the densification of these ceramics. As a consequence, their relative densities with the addition of HfB2 or HfN can be improved, and when the HfB2 and HfN contents are 30 wt %, the optimal relative densities of the TiB2-HfB2 and TiB2-HfN ceramics are 99.0% ± 0.2% and 99.4% ± 0.3%, respectively.   As can be seen, their relative densities increase with increasing HfB 2 and HfN contents from 10 wt % to 30 wt %, respectively. The relative density increments of the TiB 2 -HfN ceramics are smaller, and the relative density variation curve is relatively flat, while the relative density variation curve of TiB 2 -HfB 2 shows a bigger increment in relative density at first, and then finally shows a smaller increment. These results are ascribed to the pore number reduction with increasing the additive content, to some extent, and is derived from the higher sintering pressure (30 MPa) and the metal phases (Ni and Mo) that can efficiently reduce the sintering temperature and can accelerate the densification of these ceramics. As a consequence, their relative densities with the addition of HfB 2 or HfN can be improved, and when the HfB 2 and HfN contents are 30 wt %, the optimal relative densities of the TiB 2 -HfB 2 and TiB 2 -HfN ceramics are 99.0% ± 0.2% and 99.4% ± 0.3%, respectively. that the HfB2 additive can not only inhibit the growth of the TiB2 grains, but can also change the microstructure of TiB2-based ceramic, and that the HfN additive cannot inhibit the TiB2 grain growth.  Figure 5 presents the relative densities of the TiB2-HfB2 and TiB2-HfN ceramic tool materials. As can be seen, their relative densities increase with increasing HfB2 and HfN contents from 10 wt % to 30 wt %, respectively. The relative density increments of the TiB2-HfN ceramics are smaller, and the relative density variation curve is relatively flat, while the relative density variation curve of TiB2-HfB2 shows a bigger increment in relative density at first, and then finally shows a smaller increment. These results are ascribed to the pore number reduction with increasing the additive content, to some extent, and is derived from the higher sintering pressure (30 MPa) and the metal phases (Ni and Mo) that can efficiently reduce the sintering temperature and can accelerate the densification of these ceramics. As a consequence, their relative densities with the addition of HfB2 or HfN can be improved, and when the HfB2 and HfN contents are 30 wt %, the optimal relative densities of the TiB2-HfB2 and TiB2-HfN ceramics are 99.0% ± 0.2% and 99.4% ± 0.3%, respectively.   Figure 6 exhibits the variation of the mechanical properties of the TiB 2 -HfB 2 ceramics with changes of the HfB 2 content and variation of the mechanical properties of the TiB 2 -HfN ceramics with changes of the HfN content. In Figure 6a, with the increase of the HfB 2 content from 10 wt % to 30 wt %, the flexural strength increases from 680.49 ± 15 MPa to 708.71 ± 18 MPa; Vickers hardness increases from 19.15 ± 0.21 GPa to 21.52 ± 0.24 GPa; however, the fracture toughness decreases from 6.92 ± 0.18 MPa·m 1/2 to 5.53 ± 0.18 MPa·m 1/2 . The TiB 2 -30 wt %HfB 2 ceramic tool material exhibits better mechanical properties including flexural strength of 708.71 ± 18 MPa, which is higher than 533 MPa (the flexural strength of TiB 2 -TaC ceramics [9]), Vickers hardness of 21.52 ± 0.24 GPa that is higher than 19.8 ± 0.6 GPa (Vickers hardness of the TiB 2 -SiC-CNTs ceramics [8]), and fracture toughness of 5.53 ± 0.18 MPa·m 1/2 that is higher than 5.2 MPa·m 1/2 (fracture toughness of the TiB 2 -SiC ceramics [31]). The improvement of flexural strength and Vickers hardness is due to the relatively fine microstructure, which is in agreement with the result that the fine microstructure can improve the mechanical properties of ceramic composite materials [32]. As the HfB 2 content increases, the fracture toughness decreases gradually, which can be ascribed to the increase of the brittle rim phase. The reason is that the rim phase is mainly the complex solid solution of TiB 2 and HfB 2 , which may be a brittle phase; moreover, the wettability between HfB 2 and Ni (the wettability angle:~99 • ) is poor.  Figure 6 exhibits the variation of the mechanical properties of the TiB2-HfB2 ceramics with changes of the HfB2 content and variation of the mechanical properties of the TiB2-HfN ceramics with changes of the HfN content. In Figure 6a, with the increase of the HfB2 content from 10 wt % to 30 wt %, the flexural strength increases from 680.49 ± 15 MPa to 708.71 ± 18 MPa; Vickers hardness increases from 19.15 ± 0.21 GPa to 21.52 ± 0.24 GPa; however, the fracture toughness decreases from 6.92 ± 0.18 MPa·m 1/2 to 5.53 ± 0.18 MPa·m 1/2 . The TiB2-30 wt %HfB2 ceramic tool material exhibits better mechanical properties including flexural strength of 708.71 ± 18 MPa, which is higher than 533 MPa (the flexural strength of TiB2-TaC ceramics [9]), Vickers hardness of 21.52 ± 0.24 GPa that is higher than 19.8 ± 0.6 GPa (Vickers hardness of the TiB2-SiC-CNTs ceramics [8]), and fracture toughness of 5.53 ± 0.18 MPa·m 1/2 that is higher than 5.2 MPa·m 1/2 (fracture toughness of the TiB2-SiC ceramics [31]). The improvement of flexural strength and Vickers hardness is due to the relatively fine microstructure, which is in agreement with the result that the fine microstructure can improve the mechanical properties of ceramic composite materials [32]. As the HfB2 content increases, the fracture toughness decreases gradually, which can be ascribed to the increase of the brittle rim phase. The reason is that the rim phase is mainly the complex solid solution of TiB2 and HfB2, which may be a brittle phase; moreover, the wettability between HfB2 and Ni (the wettability angle: ~99°) is poor.  6.32 ± 0.16 MPa·m 1/2 to 7.52 ± 0.17 MPa·m 1/2 . The TiB 2 -10 wt % HfN ceramic tool material shows better mechanical properties, including flexural strength of 813.69 ± 21 MPa that is higher than 705 MPa (flexural strength of the TiB 2 -10 wt % SiC ceramics [6]), Vickers hardness of 22.59 ± 0.24 GPa which is higher than 21.85 GPa (Vickers hardness of the TiB 2 -TiC-10 wt % Ni ceramics [33]), and fracture toughness of 6.32 ± 0.16 MPa·m 1/2 that is higher than 6 MPa·m 1/2 (fracture toughness of the TiB 2 -2.5 wt % MoSi ceramics [15]). The decrease of the flexural strength and Vickers hardness is due to the increase of the defects such as the TiB 2 coarse grain and HfN agglomeration; this indicates that the defects have more negative effects on the flexural strength and Vickers hardness than the core-rim structure, although the core-rim structure is advantageous for improving the mechanical properties. The enhancement of fracture toughness is mainly attributed to the decrease of the pore number and the increase of the rim phase and TiB 2 coarse grain; decreasing the pore formation can keep the cracks from growing, which will improve fracture toughness; the rim phase of TiB 2 -HfN ceramics exhibits a higher grain boundary strength than the rim phase of TiB 2 -HfB 2 ceramics, which will provide a larger grain growth resistance for enhancing fracture toughness; in addition, TiB 2 coarse grains can consume more fracture energy in the fracturing process even though the TiB 2 is a brittle phase, which leads to the improvement of the fracture toughness.

Mechanical Properties
In order to further analyze the toughening mechanisms of TiB 2 -HfB 2 and TiB 2 -HfN ceramics, the crack propagation paths are shown in Figure 7. As can be seen, the crack propagation path in Figure 7a is straighter than that in Figure 7b; the crack deflection in Figure 7b is more obvious than that in Figure 7a; crack bridging and transgranular fracture play an important role in Figure 7a, while crack deflection, crack bridging, and transgranular fracture occupy important positions in Figure 7b, which are advantageous for enhancing fracture toughness and are the main toughening mechanisms of these ceramics. Much fracture energy will be consumed by crack bridging because crack bridging as well as crack deflection can change the direction of crack propagation (see the red circles in Figure 7), which is advantageous for improving fracture toughness. Usually the formation of the rim phase is propitious to the enhancement of fracture toughness, but in Figure 7a the rim phase shows a brittle characteristic leading to lower fracture toughness with increasing HfB 2 content as mentioned above; moreover, the relatively straight crack crossing the rim phase and TiB 2 grain will consume less fracture energy, which is harmful to the improvement of fracture toughness. However, intergranular fracture and transgranular fracture coexisted in Figure 7b, where the crack path is full of twists and turns which is advantageous to enhancing fracture toughness. In Figure 6b, with the increase of the HfN content from 10 wt % to 30 wt %, the flexural strength decreases from 813.69 ± 21 MPa to 716.37±23 MPa; the Vickers hardness decreases from 22.59 ± 0.24 GPa to 19.23 ± 0.23 GPa; however, the fracture toughness increases from 6.32 ± 0.16 MPa·m 1/2 to 7.52 ± 0.17 MPa·m 1/2 . The TiB2-10 wt % HfN ceramic tool material shows better mechanical properties, including flexural strength of 813.69 ± 21 MPa that is higher than 705 MPa (flexural strength of the TiB2-10 wt % SiC ceramics [6]), Vickers hardness of 22.59 ± 0.24 GPa which is higher than 21.85 GPa (Vickers hardness of the TiB2-TiC-10 wt % Ni ceramics [33]), and fracture toughness of 6.32 ± 0.16 MPa·m 1/2 that is higher than 6 MPa·m 1/2 (fracture toughness of the TiB2-2.5 wt % MoSi ceramics [15]). The decrease of the flexural strength and Vickers hardness is due to the increase of the defects such as the TiB2 coarse grain and HfN agglomeration; this indicates that the defects have more negative effects on the flexural strength and Vickers hardness than the core-rim structure, although the core-rim structure is advantageous for improving the mechanical properties. The enhancement of fracture toughness is mainly attributed to the decrease of the pore number and the increase of the rim phase and TiB2 coarse grain; decreasing the pore formation can keep the cracks from growing, which will improve fracture toughness; the rim phase of TiB2-HfN ceramics exhibits a higher grain boundary strength than the rim phase of TiB2-HfB2 ceramics, which will provide a larger grain growth resistance for enhancing fracture toughness; in addition, TiB2 coarse grains can consume more fracture energy in the fracturing process even though the TiB2 is a brittle phase, which leads to the improvement of the fracture toughness. In order to further analyze the toughening mechanisms of TiB2-HfB2 and TiB2-HfN ceramics, the crack propagation paths are shown in Figure 7. As can be seen, the crack propagation path in Figure 7a is straighter than that in Figure 7b; the crack deflection in Figure 7b is more obvious than that in Figure 7a; crack bridging and transgranular fracture play an important role in Figure 7a, while crack deflection, crack bridging, and transgranular fracture occupy important positions in Figure 7b, which are advantageous for enhancing fracture toughness and are the main toughening mechanisms of these ceramics. Much fracture energy will be consumed by crack bridging because crack bridging as well as crack deflection can change the direction of crack propagation (see the red

Conclusions
TiB 2 -based ceramic tool materials reinforced by HfB 2 and HfN additives have been fabricated by hot pressed sintering. The effects of HfB 2 and HfN additions on their microstructures and mechanical properties were investigated. The results showed that the HfB 2 additive can inhibit the TiB 2 grain growth and can change the morphology of some of the TiB 2 grains from bigger polygons to smaller polygons or longer ovals, which is favorable for the formation of a relatively fine microstructure, while the HfN additive tends to agglomerate. With increasing HfB 2 and HfN contents from 10 wt % to 30 wt %, the relative densities of these ceramics increased gradually. The relatively fine microstructure improved the flexural strength and Vickers hardness of the TiB 2 -HfB 2 ceramics. The poor wettability between HfB 2 and Ni resulted in the formation of weak grain boundary strength and the complex solid solution of TiB 2 -HfB 2 is a brittle phase, which led to the decrease of fracture toughness of the TiB 2 -HfB 2 ceramics. The increase of the defects such as the TiB 2 coarse grain and HfN agglomeration resulted in the decrease of the flexural strength and Vickers hardness of the TiB 2 -HfN ceramics; the decrease of the pore number and the increase of the rim phase and TiB 2 coarse grain are advantageous for the enhancement of fracture toughness. The toughening mechanisms of the TiB 2 -HfB 2 ceramics mainly included crack bridging and transgranular fracture, while the toughening mechanisms of the TiB 2 -HfN ceramics mainly included crack deflection, crack bridging, transgranular fracture, and the core-rim structure. The TiB 2 -30 wt % HfB 2 ceramic tool material exhibited better mechanical properties including a flexural strength of 708.71 ± 18 MPa, Vickers hardness of 21.52 ± 0.24 GPa, and fracture toughness of 5.53 ± 0.18 MPa·m 1/2 . The TiB 2 -10 wt % HfN ceramic tool material showed better mechanical properties including a flexural strength of 813.69 ± 21 MPa, Vickers hardness of 22.59 ± 0.24 GPa, and fracture toughness of 6.32 ± 0.16 MPa·m 1/2 .