Microstructure Evolution and Mechanical Properties of B4C-Reinforced TC11 + xFe Composites Fabricated by HIP

Abstract

The present study involved (TiB + TiC)/TC11 (Ti-6.5Al-3.5Mo-1.2Zr-0.3Si) + xFe titanium matrix composites (TMCs) reinforced by in situ TiB whiskers and TiC particles fabricated by hot isostatic pressing. Microstructure observation reveals a substantial distribution of in situ reinforcements, which form a network-reinforced structure at the prior particle boundaries of the TC11 matrix. The micro–nanoscale TiB whiskers and TiC particles within and surrounding this network serve as effective dislocation pinning. The enhancement of mechanical properties can be attributed to load-bearing strengthening, fine-grain strengthening, and dislocation strengthening. The hardness and compressive strengths were investigated through mechanical properties testing. The hardness increased by 19.4% (2 wt% B₄C-reinforced composites) compared with TC11 alloy. However, the addition of 2 wt% Fe at the same B₄C level (2 wt% B₄C + 2 wt% Fe-reinforced composites) resulted in a significant increase in hardness by 37.5% and 15.2% in compressive strengths of TMC and can be attributed to the solid solution strengthening effect and higher dislocation density provided by the addition of Fe. In addition, the optimal overall properties can be achieved by strictly regulating the addition ratio of 2 wt% Fe and 1 wt% B₄C, allowing for a compressive strength of 2301 MPa while still maintaining a compressive strain of 24.6%.

1. Introduction

Titanium alloys are widely used in the petrochemical, biomedical, and aerospace fields due to their high strength, low density, and corrosion resistance [1,2,3,4]. Due to rapid advancements in various fields, the properties requirements for titanium alloys are becoming increasingly demanding. Titanium matrix composites (TMC) have emerged as a pivotal solution to tackle these challenges due to their exceptional properties, such as high specific strength and temperature stability [5,6,7]. In the latter quarter of the 20th century, TMC was extensively researched in the United States and Europe. The incorporation of ceramic reinforcements, such as TiC, SiC, and B₄C, is widely regarded as an effective method to achieve superior properties [8,9]. Reinforcement with different compositions can be generated through in situ reactions using the principles of powder metallurgy (PM), resulting in excellent metallurgical interfacial bonding and thermodynamic stability [10,11]. This process, such as hot isostatic pressing (HIP), significantly enhances the mechanical properties at both room temperature and high temperatures [12].

Solving the brittleness of TMCs has been the focus of extensive attention by researchers. The mechanical properties of the composite material primarily depend on the microstructure and properties of the matrix, as well as the type, mass fraction, distribution, and size of the reinforcement material. Additionally, they are influenced by the extent of interfacial bonding between the matrix and reinforcement. Researchers have improved both strength and plasticity by constructing a network reinforcement structure. Liu et al. [11] enhanced both the strength and ductility by preparing in situ reinforced (TiB + TiC)/Ti-6Al-4V composites with a network structure. Zhong et al. [13] constructed a two-scale network structure of TiB nanowires and micro-TiC particles combined with needle-like silicide, overcoming the strength–extensibility paradox. The in situ route of forming TMCs is a highly efficient and cost-effective method with a clean and strong reinforcement–matrix interface bonding, and the reinforcement is frequently precipitated from the matrix as either whisker or particle reinforcements [14,15]. Moreover, researchers have enhanced the ratio of TiB whiskers and formed nano-TiB whiskers to tackle the strength–ductility paradox [16,17]. Therefore, we introduced hybrid in situ TiB whisker and TiC particles through the process of PM-HIP, which has been proved effective in strengthening materials [18]. Furthermore, the addition of nano-B₄C powder leads to the formation of reinforcement with micro and nano sizes, and nano-sized reinforcements can provide significant strengthening in TMCs [19,20].

Moreover, the addition of Fe plays a crucial role in optimizing the strength and plasticity matching of titanium alloys. Researchers confirmed that Fe can affect the lattice distortion of titanium alloys, leading to significant solid solution strengthening [21,22,23]. Wang et al. [24] prepared in situ alloyed Ti–5Fe using selective laser melting (SLM), which achieved high tensile strength (865 MPa) with good elongation (12%). Niu et al. [25] studied Ti-Fe alloy and proved that the best overall properties were achieved by adding 2 wt% Fe.

Despite the significant enhancement of mechanical properties in titanium alloys through Fe addition, there is a lack of reports on further enhancing TMCs by the addition of Fe, and its influence on the mechanical properties and microstructure of TMCs remains unreported. Therefore, the addition of 1, 2 wt% Fe and addition of 0, 1, 2, 5 wt% B₄C are selected to enhance the Ti-6.5Al-3.5Mo-1.2Zr-0.3Si (TC11) alloy. In this study, the (TiB + TiC)/TC11 composite with a network-reinforced structure was prepared in situ by HIP by adding nano-B₄C powder to the TC11 matrix. In addition, Fe is added to obtain the best properties by regulating the content of B₄C and Fe. The microstructure and properties of all TMCs were studied and compared, providing a reference for the study of TMCs with good performance.

2. Materials and Methods

2.1. Material and Manufacture Process

The experimental material used in this work was titanium matrix composite prepared by two-step hot isostatic pressing (HIP). Pre-alloyed α + β titanium alloy Ti-6.5Al-3.5Mo-1.2Zr-0.3Si (referred to as TC11 hereafter) powder fabricated by a plasma rotating electrode process (PREP) (SHANGJI Co., Nanjing, China) with particle size ranges from 53 μm to 106 μm was used as a matrix material. The TC11 alloy was reinforced with in situ reinforcements through the incorporation of nano-B₄C powder with an average particle size of 50 nm. Additionally, nano-Fe powder with an average particle size of 50 nm was also added to investigate the microstructure and mechanical properties of the TMCs with the addition of Fe. The proportions of the three powders are shown in Table 1. The three powders were placed in a vacuum drying oven and heated at 100 °C for 4 h before being put into a planetary ball-milling machine. The low-energy milling processes were conducted at 200 rpm for 6 h with an agate ball-to-powder weight ratio of 4:1. The ball-milled powders were loaded into stainless steel capsules with an inner diameter of 40 mm, a length of 150 mm, and a wall thickness of 2.5 mm, followed by 6 h of degassing at a vacuum of 1 × 10⁻³ pa and a temperature of 600 °C. The HIP process was conducted at 1050 °C/150 MPa for 4 h. After being cooled to room temperature, stainless steel capsules were removed by machining, followed by a two-step HIP sintering at 1150 °C/150 MPa for 4 h. The flow chart illustrating the fabrication process of (TiB + TiC)/TC11 composites through the HIP technique is presented in Figure 1.

2.2. Microstructure and Morphology Characterization

The samples were separated from the prepared TMCs by wire cutting. The specific dimensions of each sample are depicted in Figure 2. The samples were subjected to X-ray diffraction (XRD, Rigaku SmartLab SE, AXT, Tokyo, Japan) for phase identification, operated with a Cu target at 35 kV and 40 mA using a continuous scan mode. A scan was conducted at 10°/min over a wide range of 30–80°. The samples for scanning electron microscope (SEM) observation were ground, polished, and etched by Kroll’s reagent (10 vol% HF, 20 vol% HNO₃, and 70 vol% H₂O). Microstructure and morphology characterizations were conducted through a SEM (JSM6490LV, JEOL, Osaka, Jpapn). The microstructure of initial powders was observed using a SIGMA field emission SEM (Zeiss, Oberkochen, Germany). The specimens were polished by argon ion-beam polishing (EM TIC 3X, Leica, Solms, Germany) at a voltage of 6 KV for 120 min and 4 KV for 30 min before electron backscatter diffraction (EBSD, C-SWIFT, Oxford, UK) observation. EBSD and AZtec software (Aztec Crystal 2.1.2, Oxford Instruments, Oxford, UK) were adopted to analyze the orientation of α-Ti, β-Ti, TiB, and TiC phases and phase content.

2.3. Mechanical Properties Testing

According to the Chinese Standard GB/T 7314-2017 [26], the room-temperature compression tests were executed on a universal testing machine (Instron 5569, Canton, MA, USA) with a loading rate of 0.1 mm/min, using a compression specimen size of Φ4 mm × 6 mm as shown in Figure 2.

Vickers hardness (HV) was measured using an HVS-1000 Vickers Hardness Tester, with a force of 20 N applied for 15 s. All the samples for hardness measurement underwent a sequential polishing process using sandpaper of varying grades, including 200, 400, 800, 1200, 1500, and 2000. A minimum of 15 tests were conducted to obtain the average result. The position for each test was set at a parallel interval of 20 μm, with the center of the sample as the starting point.

3. Results

3.1. Microstructure Characteristics of the Initial Powders

The microstructure characteristics of the B₄C powder, Fe powder, pre-alloyed TC11 powder, and mixed powder are summarized in Figure 3. Both the pre-alloyed TC11 powder (Figure 3a) and mixed powder (Figure 3d) exhibit excellent sphericity with no discernible defects, resulting in desirable powder flowability. This ensures the powder is well vibrated and fills the stainless steel capsules, which is conducive to the densification process [27]. The nano-B₄C powder (Figure 3b) and nano-Fe powder (Figure 3c) are uniformly adhered to the TC11 pre-alloy powder particles. The morphology of a single amplified particle of mixed powder, as depicted in Figure 3d, demonstrates a uniform combination of micro–nanoparticles achieved through low-energy ball milling.

3.2. Phase Composition

The XRD patterns of TMCs and TC11 alloy fabricated by HIP are shown in Figure 4. The results show that the α-Ti, β-Ti, TiB, and TiC peaks are clearly detected without the observation of the diffraction peak for B₄C. The XRD analysis reveals a low strength of β-Ti peaks, particularly in TC11, TMC1, and TMC2 without the addition of Fe, indicating limited precipitation of β-Ti grains. The addition of the β-stable element Fe results in the identification of more pronounced peaks in TMC3-F, TMC4-F, TMC5-F, and TMC6-F, as shown in the amplified spectrum B (Figure 4c). The raw material B₄C had been converted to the TiB whisker (TiB_w) and TiC particle (TiC_p) phases through the in situ reaction as 5Ti + B₄C→4TiB + TiC during the HIP process [28,29,30]. The intensity of TiC and TiB peaks increases with further addition of B₄C as shown by the amplified spectrum A in Figure 4b. The position of the matrix diffraction peak also shifted to the left due to an increase in the lattice constant of the TC11 matrix resulting from the formation of an interstitial solid solution with a solid solution of B and C atoms [31]. The HIP process did not introduce more reactions.

3.3. Morphology of Microstructure

The SEM images of TC11 alloy and B₄C-reinforced TC11 composites demonstrate the evolution of microstructure, as shown in Figure 5 and Figure 6. As shown in Figure 5a, the microstructures of TC11 matrix alloy consist of equiaxial α phase, platelet α phase, and residual β phase (β_r). The continuous grain boundary α phase (α_GB) and α/β colony structure with different orientations are also observed. In Figure 5b,c, it is observed that the in situ TiB and TiC reinforcements primarily accumulate at the prior particle boundaries (PPBs) and form a network structure. In addition, traces of nano-TiB whiskers (TiB_w) and nano-TiC particles (TiC_p) can be found in the TC11 matrix slightly away from the PPBs, indicating that C atoms and B atoms in the (TiB + TiC)-rich region diffused a certain distance into the interior of the network structure. With the increase in B₄C addition, TiB_w and TiC_p grow and agglomerate on PPBs, resulting in the formation of a coarser and more continuous network structure. This alteration adversely affects the mechanical properties of TMCs [32]. It is noteworthy that the incorporation of 5 wt% B₄C and insufficient solid phase sintering resulted in the presence of pores at the PPBs in TMC6, with a nearly circular shape and an average size of 6.5 μm.

Figure 6 presents the microstructures of the B₄C-reinforced TC11 composites with the addition of Fe. Compared with TC11 alloy, the microstructures of TMCs tend to be more equiaxial, which indicates that (TiB + TiC) dual phases can facilitate the formation of equiaxed grains. The comparison of TMC3-F, TMC5-F, and TMC6-F reveals that with the same amount of added Fe, an increase in added B₄C from 1 wt% to 3 wt% results in a higher reinforcement content and a more continuous network structure. When comparing TMC1, TMC4-F, and TMC5-F, it can be observed that an increase in added Fe from 0 wt% and 1 wt% to 2 wt% leads to more nano-TiB_w (TiB whiskers) clusters with diverse orientations and more nano-TiC_p (TiC particles) stacked densely between TiB_w near the network structure when the same amount of B₄C is added. In addition, more micron-scale TiB_w and TiC_p precipitate from the matrix. These changes prove that the addition of Fe promoted the in situ reaction and the distribution of elements. The addition of Fe can slightly alter the phase content and size of the matrix, making it easier to form a narrower α phase, while the width and content of the β phase may increase accordingly.

The EDS mapping was introduced to determine the chemical composition of microstructure and reinforcing phases in (TiB + TiC)/TC11 composites fabricated by HIP. To determine the distribution of Fe elements in the composite, EDS scanning was conducted using TMC5-F as a representative sample, as shown in Figure 7a. The brighter regions indicate a higher concentration of these elements. EDS point scanning was also carried out to further verify the composition of the reinforcement-rich region, the results of which are consistent with the XRD results in Figure 4, indicating that the reinforcement consists of TiB and TiC. It is obvious that the matrix is primarily composed of Ti and Al, which are the most abundant elements in its composition. The B and C elements uniformly dissolve in the matrix and are enriched at the prior particle boundaries (PPbs), where they undergo in situ reactions with Ti elements to form the (TiB + TiC) network structure. The in situ nano-TiB_w exhibit a high aspect ratio, with a diameter of approximately 100 nm and lengths starting from several microns. They are arranged in parallel clusters and distributed around the network structure. TiC_p and TiB_w are interwoven and distributed near the network structure, with a small amount diffused into the matrix. The multiple reinforcements contribute to enhancing the properties of TMCs [33]. In addition, the Fe element is dissolved in the matrix, particularly in regions where the β phases are present, which has a distinctive performance in solid solution strengthening and improves the working properties [34,35].

3.4. EBSD Analysis

EBSD analysis was used to further investigate the crystal orientation and phase characteristics of 2 wt% B₄C-reinforced TC11 composites with added 2 wt% Fe or not. Figure 8 shows the band contrast (BC), inverse pole figure (IPF), and phases maps of the HIPed TMC1 and TMC5-F. The results show that the primary detection phase of TMC1 and TMC5-F was α-Ti with a relatively low content of β-Ti. According to the phase maps (Figure 8c,f), the distribution of the TiB and TiC reinforcement-rich region is marked in Figure 8a,d. Figure 8b indicates that the α-Ti grain orientations of the HIPed sample were randomized with {$\overline{1}$2$\overline{1}$0} and {01$\overline{1}$0}. After adding Fe, a significant transformation from {0001} and {01$\overline{1}$0} crystal orientations to {$\overline{1}$2$\overline{1}$0} was observed, as shown in Figure 8e. This is because the addition of Fe exhibits distinct selectivity towards the grain orientation of the α phase and, as Fe addition increases, the α phase grain orientation becomes more concentrated, leading to a more pronounced isotropy [36]. The high-angle grain boundaries (>10°) predominated with a high proportion of 73.7% (TMC1) and 70.0% (TMC5-F), suggesting a rise in the number of small-angle grain boundaries when Fe is added. After HIP, the majority of β-Ti phase regions in TMC5-F were undetectable, consistent with the findings obtained from XRD analysis.

Furthermore, the effect of the addition of Fe on the texture strength of the α phase in TMC1 and TMC5-F was further analyzed, as shown by pole figure diagrams in Figure 9a,b. Both TMC1 and TMC5-F demonstrate greater texture strength along the (0001) orientation [37]. It is noteworthy that TMC5-F exhibits greater texture strength in the (11−20) and (10−10) directions compared to TMC1 due to the addition of 2 wt% Fe, resulting in a preferred orientation in the texture of the α phase compared to TMC1. In total, the HIPed TMC has no obvious texture, so its mechanical properties exhibit isotropy [27].

In order to qualitatively and quantitatively describe and compare the dislocation density of B₄C-reinforced TC11 composites with and without the addition of Fe, KAM maps were used to determine the change trend of dislocation density. An equation was introduced to determine the geometrically necessary dislocations (GNDs) for TMC1 and TMC5-F. The equation is as follows [38,39,40]:

(1)${\rho }^{GND}={\displaystyle \frac{2{KAM}_{ave}}{ub}}$

Equation (1): ${\rho }^{GND}$ is the GND density of the measurement point; ${KAM}_{ave}$ is the average value of KAM derived from EBSD software; the EBSD scan step size is 180 nm, denoted as $u$; the Burgers vector of the titanium alloy, which is 2.951 nm for HCP titanium, is denoted as $b$. Figure 10 presents the KAM distribution maps and the computed GND density for TMC1 and TMC5-F. The value of the GND of the 2 wt% B₄C-reinforced TC11 composites (TMC1) is calculated as 1.89 × 10¹⁴ m⁻² and increases to 2.09 × 10¹⁴ m⁻² (increased by 10.6%) after the addition of 2 wt% Fe (TMC5-F). Compared with TMC1, TMC5-F has a higher dislocation density which can be attributed to the addition of Fe. Fe element can induce the precipitation of the β phase, resulting in more α/β interfaces. In addition, the addition of Fe promotes the precipitation of in situ reinforcement, and the interface number of reinforcements increases, resulting in a higher dislocation density. High dislocation density could significantly enhance tensile strength [41].

3.5. Room Temperature Uniaxial Compression Experiment

Figure 11a demonstrates the compressive stress–strain curves of 2 wt% and 5 wt% B₄C-reinforced TC11 composites. With the increase in the addition of B₄C, the compressive yield strength reaches 1327 MPa (TMC1) and 1586 MPa (TMC2), respectively, which increased by 30.1% and 55.5% compared with TC11 alloy (1020 MPa). The stress–strain curves of 2 wt% B₄C-reinforced TC11 composites with the addition of 0, 1, 2 wt% Fe are shown in Figure 11b. The addition of Fe is observed to significantly enhance both the compressive yield strength and compressive strength of the material, resulting in a 2.8% increase in compressive strength for TMC4-F and an 8.9% increase for TMC5-F compared to TC11. Moreover, when compared to TMC1 without added Fe, TMC4-F exhibits a 6.8% increase in compressive strength while TMC5-F demonstrates a remarkable 13.2% improvement. This indicates that the addition of 2 wt% Fe can improve the compressive strength of the material and maintain a certain plasticity. Therefore, the compression property of 1, 2, 3 wt% B₄C-reinforced TC11 composites with the addition of 2 wt% Fe is studied as shown in Figure 11c. It can be found that the compressive properties of the composites do not exhibit improvement with an increase in B₄C content when Fe is at 2 wt%. Moreover, the composites containing 1 wt% B₄C and 2 wt% (TMC3-F) Fe demonstrate superior compressive properties. TMC3-F has the highest compressive strength, reaching 2301 MPa, which represents a 15.2% improvement compared to TC11, while maintaining a certain compressive strain (24.6%).

The compressive yield strength of all the TMCs markedly exceeds that of the TC11 alloy, as can be seen from the box selection area in Figure 11d. In general, the compressive strength of TMCs increased first and then decreased with the increased addition of B₄C. The comparison of TMC1, TMC4-F, and TMC5-F indicates that the addition of Fe can substantially increase the compressive strength without compromising the compressive strain. The addition of 2wt% Fe is observed to yield the most effective strengthening effect. In this case, the addition of 1wt% B₄C achieves optimal strength enhancement. Consequently, TMC3-F demonstrates the best balance between strength and plasticity. It is noteworthy that the strains of TMC2 and TMC6-F exhibit a significant reduction, despite their high yield strength. This indicates brittle fracture characteristics as both compressive strength and strain are considerably reduced.

Moreover, the hardness of all the TMCs is significantly increased and the increasing trend in hardness is consistent with the compressive yield strength, as shown in Figure 11e. The comparison of TC11, TMC1, and TMC2 reveals that (TiB + TiC) reinforcements can significantly harden the TC11 matrix, and the addition of more B₄C led to higher hardness. The addition of B₄C at 2 wt% and 5 wt% led to corresponding increases in the composite hardness by 19.4% and 34.5%, respectively. Comparison of TMC1, TMC4-F, and TMC5-F reveals that the addition of Fe can significantly enhance the hardness. The hardness of TMC4-F is 15.2% higher than that of TMC1 and 37.5% higher than that of the TC11 alloy due to the addition of 1 wt% Fe. The addition of a reasonable amount of reinforcement and the appropriate adjustment of Fe have been shown to achieve better performance improvement in TMCs, as evidenced by their excellent compressive strength, hardness, and high strain, surpassing the simple addition of B₄C.

4. Discussion

Figure 12 shows the schematic microstructure diagram of microstructural evolution of the (TiB + TiC)-reinforced composites. TiB whiskers and TiC particles in the micro- and nanoscale were formed in the TC11 matrix of the HIPed TMCs, indicating a solid–solid reaction during the sintering of B₄C and Ti powders to form in situ TiC and TiB. After low-energy ball milling, the B₄C particles were uniformly distributed on the TC11 matrix particles, which provided the carbon source and boron source. The formation of TiB and TiC phases resulted from the reaction between B₄C and Ti as follows:

5Ti + B₄C = 4TiB + TiC(2)

Under the high sintering temperature of 1050 °C and the isostatic pressure of 150 MPa, the rapid diffusion of titanium, carbon, and boron atoms was facilitated by the intense thermal atomic motion, resulting in nucleation and growth of TiB_w and TiC_p. The eutectic transformation between titanium and iron takes place at a temperature of 1085 °C; therefore, the sintering temperature is strictly controlled at 1050 °C to prevent reaction between the stainless steel capsules and the TC11 matrix [42]. Furthermore, by implementing a two-step HIP process after the removal of the stainless steel capsule, it ensured that the reinforcements were fully formed through in situ reactions. The B₄C particles tend to aggregate at prior particle boundaries (PPBs), which provide abundant sources of boron and carbon. As a result, during the solid phase reaction, larger micron-sized TiC_p and TiB_w appear at PPBs. Furthermore, nanoscale TiB_w and TiC_p were observed within the matrix at a specific distance from the PPBs as a result of diffusion of boron and carbon atoms into the interior of the network structure [43]. Thus, the concentration of C and B elements mainly contributes to the grain size of TiB_w and TiC_p, while the growth morphologies in the reinforcements primarily contribute to their crystal structures. The growth of TiB_w is oriented along the [010] crystallographic axis in the longitudinal direction, resulting in TiB_w exhibiting a needle-like morphology [44,45]. The TiC grows in each crystal planes and tends to form granules.

The compression test and hardness test results indicate that the addition of the B₄C powder to the TC11 matrix improved the mechanical properties, and the addition of Fe provides an additional strengthening effect. The network structure, featuring the distribution of in situ TiB and TiC, effectively transfers the load while preserving a certain level of matrix plasticity. Microstructure formation shows that in situ networks refined the matrix grain, leading to a great improvement in mechanical properties. The nanoprecipitated phases TiC and TiB within the network simultaneously contribute to load transfer and precipitated phase strengthening. Since the interfacial strength of TiB-Ti and TiC-Ti is much higher than that of Ti-Ti, the formation of a large number of nano-TiB_w and nano-TiC_p provides more TiB-Ti and TiC-Ti interfaces, which leads to load-bearing strengthening effect while providing fine-grain strengthening, significantly increasing the strength of the material [46,47]. The solid solution strengthening caused by interstitial B and C elements is theoretically limited and can be disregarded. Since the thermal expansion coefficients of TiB and Ti are similar, the dislocation strengthening caused by TiB can be ignored [48]. The presence of TiC_p at the grain boundaries can account for the dislocation pinning effect and contribute to the enhanced strength observed in TMCs. The deformation compatibility between Ti and TiC differs, resulting in stress concentration at the interface between hard TiC particles and Ti. However, excessive addition of B₄C promotes crack propagation and defect formation, which accounts for the presence of pores in TMC2. Therefore, incorporating an appropriate amount of B₄C to induce nanoscale in situ reinforcements can mitigate this effect and provide dislocation strengthening. The reinforcements act as obstacles against dislocations, thereby enhancing the mechanical properties of materials.

The additional strengthening effect provided by Fe is due to the solid solution strengthening of Fe. Moreover, the addition of Fe can further facilitate the in situ formation of TiB_w and TiC_p, as shown in Figure 13. As shown in Figure 4b, the presence of the TiFe-base intermetallic phase was not detected in TMC, and Fe was incorporated into TMCs as a β stable element, resulting in an increased intensity of the β peak attributed to the addition of Fe. In addition, Fe exhibits a significant solid solution strengthening effect due to its strong ability to induce lattice distortion in titanium alloy. The strain caused by the lattice distortion can increase the Young’s modulus and bulk modulus of the β phase, thus affecting the plasticity of the material [49]. Thus, the combination of in situ reinforcement and Fe solid solution strengthening can result in an additional strengthening effect. However, S. Pouzet et al. [50] reported that the content of B₄C is typically limited to no more than 3 wt% in order to ensure sufficient plasticity which is consistent with our experimental results. Therefore, the precise control of the reinforcement and Fe content in the particle is crucial for achieving optimal performance balance.

5. Conclusions

In this work the microstructure evolution and mechanical properties of (TiB + TiC)/TC11 + xFe composites fabricated by HIP were systematically investigated. The findings can be summarized as follows:

• (1). After low-energy ball milling, the nano-B₄C and nano-Fe powders were uniformly attached to the micron-sized TC11 particles. The HIP sintering process facilitated the formation of in situ TiB and TiC reinforcements at the PPBs, resulting in well-defined network-reinforced structures.

• (2). Nano-TiB_w with high length-to-diameter ratio and granular TiC_p are distributed around the network structure, while a small amount of TiC_p and TiB_w is also dispersed within the network structure due to element diffusion, which provided load-bearing strengthening, fine-grain strengthening, and dislocation strengthening.

• (3). The B₄C-reinforced TC11 composites presented an increment in hardness and compressive yield strength, which gradually increased with more in situ TiB and TiC formation. After the addition of Fe, the composite demonstrates enhanced compressive strength, compressive strain, and hardness.

• (4). The addition of Fe can provide additional solid solution strengthening, thereby further enhancing the strength while maintaining compressive strain. Furthermore, by optimizing the ratio of B₄C and Fe additions, superior comprehensive performance can be achieved with a 15.2% increase in compressive strength while preserving excellent compressive strain of 24.6% (TMC3-F).

Author Contributions

Conceptualization, L.Z., H.C., L.F. and Y.S.; methodology, S.Q.; software, S.Q.; validation, S.Q., N.W., F.C., Y.S., J.Z. and H.C.; formal analysis, S.Q.; investigation, S.Q. and F.C.; resources, H.C., L.F. and J.Z.; data curation, S.Q.; writing—original draft preparation, S.Q.; writing—review and editing, S.Q. and Y.S.; supervision, S.Q., N.W., F.C., Y.S., J.Z., L.F. and L.Z.; project administration, Y.S., J.Z., H.C. and L.Z.; funding acquisition, H.C. and L.Z. All authors have read and agreed to the published version of the manuscript.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Conflicts of Interest

Author Jiong Zhao was employed by the company Jiangsu Tiangong Technology Co., Ltd. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Footnotes

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Figures and Table

Figure 1. Flow chart of (TiB + TiC)/TC11 composites fabricated by hot isostatic pressing.

Figure 2. Sample preparation for SEM, XRD, EBSD, and room-temperature compression tests.

Figure 3. Microstructure characteristics of powders: (a) TC11 powder; (b) Fe powder; (c) B4C powder; (d) mixed powder and amplified single particle.

Figure 4. (a) XRD patterns of TC11 and TMCs; (b) amplified spectrum marked as A; (c) amplified spectrum marked as B.

Figure 5. SEM micrographs of TC11 alloy and B4C-reinforced TC11 composites: (a) TC11; (b) TMC1; (c) TMC2.

Figure 6. SEM micrographs of B4C-reinforced TC11 composites with Fe addition: (a) TMC3-F; (b) TMC4-F; (c) TMC5-F; (d) TMC6-F.

Figure 7. EDS results of TC11-2wt% B4C with added 2wt% Fe (TMC5-F): (a) EDS mapping of TMC5-F; (b) EDS point scanning of (TiB + TiC)-rich region in TMC5-F.

Figure 8. EBSD characterization of TMC1 (a–c) and TMC5-F (d–f): (a,d) BC maps; (b,e) IPB maps; (c,f) phase maps.

Figure 9. Pole figure of α phase: (a) TMC1; (b) TMC5−F.

Figure 10. The KAM distribution, along with the mean GND density computed from the KAM maps: (a) TMC1; (b) TMC5−F.

Figure 11. (a) The stress–strain curves of 2 wt% and 5 wt% reinforced TC11 composites; (b) The stress–strain curves of 2 wt% B4C-reinforced TC11 composites with the addition of 0, 1, 2 wt% Fe; (c) The stress–strain curves of 1, 2, 3 wt% B4C-reinforced TC11 composites with the addition of 2 wt% Fe; (d) The stress–strain curves of TMCs and TC11 alloy formed by HIP; (e) The Vickers hardness comparison of TMCs and TC11 alloy.

Figure 12. Schematic diagram of microstructural evolution of the HIPed (TiB + TiC)/TC11.

Figure 13. Schematic diagram of microstructural evolution of the HIPed (TiB + TiC)/TC11 + xFe.

Table 1

Mixture ratios of TC11, B₄C, and Fe powder (wt.%).

Number	TC11	B4C	Fe
TC11	100	0	0
TMC1	98	2	0
TMC2	95	5	0
TMC3−F	97	1	2
TMC4−F	97	2	1
TMC5−F	96	2	2
TMC6−F	96	3	2

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