Investigating the binding properties of NbC/Fe-based composite layer and HT300 through experiments and simulations

In this work, the hot-pressing diffusion method was used for the fabrication of a novel composite structure. More specifically, by keeping the temperature constant at 1000 °C and applying a pressure value of 40 MPa for 60 min, α 90 min and 120 min, respectively, the NbC/Fe composite layer on the surface of HT300 was formed. The microstructure, element distribution, microhardness, bonding property and scratch deformation characteristics of NbC/Fe composite layer were studied, and the fracture mode was studied by simulation and tensile test at the micro level. The results show that the main components of the NbC/Fe composite layer prepared in the experiment are α- Fe and NbC, the composition of the composite layer is pure. The thickness of NbC/Fe composite layer prepared with 60 min, 90 min and 120 min holding time is 5 μm, 15 μm and 23 μm. The hardness of the composite layer can reach 2096.4 HV0.1; The bonding property between the NbC/Fe composite layer and the matrix is the best when the heat preservation is 120 min. Because the tensile fracture is brittle and the fracture location is in the NbC/Fe composite layer, the bonding strength between the composite layer and the matrix is greater than 297MPa, which has excellent bonding properties. In the scratch test, the longer the holding time is, the stronger the bonding ability between the reinforcing layer and the matrix is, 41.2N (90 min) and 75.75N (120 min) respectively. The fracture mechanism in the NbC/Fe composite layer was simulated by abaqus. The fracture of the composite layer was caused by the propagation of microcracks caused by the stress concentration at the sharp corner of square NbC particles in the layer.


Introduction
Millions of tons of steel materials are consumed by mining, metallurgy, construction and coal departments every year in China. The failure of the steel parts caused by serious damage to the surface accounts for about 70%-80% of the total damage events. Along these lines, surface enhancement is regarded as one of the effective ways to solve this problem. Gray cast iron is an important material for casting-based processes in machinery manufacturing because of its uniform and compact structure, high tensile strength and good mechanical properties [1][2][3]. It can be also used for bed guides, lathes, punching machines, beds with large forces, spindle box gears, etc. There are many ways to strengthen composite materials, including particle reinforcement [4,5], Infiltration [6,7]. Fiber reinforced [8]. At the same time, it is particularly important to select the appropriate enhancement phase, as far as the reinforcing phase of the steel matrix composites is considered, niobium carbide has the characteristics of high melting point, high hardness and modulus [9][10][11]. In addition, it exhibits good wettability with an iron matrix. Thus, it is considered an ideal reinforcing material.
At present, the preparation methods of NbC/Fe reinforced iron matrix composites include the in situ hot pressing method [12], laser cladding method [13][14][15], self-propagating high-temperature synthesis method [16,17], spark plasma sintering method [18][19][20], powder metallurgy [21,22], etc. However, the implementation of a high reaction temperature, in conjunction with high equipment requirements and poor interfacial bonding Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. are significant disadvantages of these methods. Ya-bin Cao [11] prepared a nanometer niobium carbide reinforced layer on an iron substrate by a employing laser cladding process. Due to the decomposition and reprecipitation procedures of the NbC reinforcement phase of the external particles, the content and morphology of NbC were significantly affected. Moreover, both the dissolution of Nb atoms and the precipitation of NbC can improve the microhardness and wear resistance of the coating. Hadian et al [29] prepared NbC-M48 high-speed steel composites with a volume fraction of 88% NbC by enforcing the vacuum hot pressing technique, with a hardness of 14.91 ± 0.14 GPa and a fracture toughness (K IC ) of only 1.56 ± 0.3 MPam 1/2 . Sen [23] et al prepared NbC coating on the surface of AISI 1040 steel by using thermal reaction-diffusion technology. The coating had no pores, and a uniform structure with a hardness value of 2512 ± 190HV 0.01 was extracted. The friction and wear test results show that the NbC-based coating possesses good abrasion resistance.

Experimental materials and methods
The preparation process is shown in figure 1. HT300 and high-purity niobium plate with a purity of >99.7% were used as raw materials, and the two were processed into cylinders (sheets) of Φ20 mm × 16 mm and Φ20 mm × 0.5 mm by wire EDM, and then were subjected to metallographic treatment. The specimens were stacked in the order of Nb/HT300, wrapped with graphite paper and placed within a graphite mold, with graphite indenters placed on both sides. They were placed in a rapid hot pressing sintering furnace, and kept at a temperature/pressure value of 1000°C−40 MPa in a vacuum environment for 60 min, 90 min and 120 min, respectively. To prevent gas contamination, the reaction is always kept under a vacuum of less than 0.1Pa.
The layer thickness and structure of the composite layer were characterized by using a JSM-6700F field emission scanning electron microscope and a Merlin Compact Zeiss scanning electron microscope (SEM). The Vickers hardness of the composite NbC/Fe composite layer was characterized by a DHV-1000Z digital microhardness tester, where the applied load was 0.1 kg and the pressure holding time was 10 s. The NbC-based composite layer on the ironbased surface was scratched by using an automatic scratch tester for coating adhesion. The continuous loading of 0 to 100 N was achieved at a speed of 100 N min −1 , and the scratch length was 5 mm. The acoustic emission curve was also used to determine the limit load of the reinforcement layer adhesion. A WDW-20 microcomputer-controlled electronic universal testing machine was used for the tensile experiment, and the experimental conditions were the application of uniaxial loading and stretching at room temperature and a constant strain rate of 0.45 mm min −1 .

Results and discussion
3.1. Phase composition and microstructure of NbC/Fe-based composite layer After the Nb/HT300 was kept at 1000°C−40 MPa for 60 min, 90 min and 120 min, respectively, the measured x-ray diffraction patterns of its cross-section are shown in figure 2. Compared with the standard PDF card [24] it can be seen that the phase composition of the NbC/Fe composite layer is composed of a face-centered cubic NbC structure and ferrite (α-Fe, no other impurities). It is interesting to notice that when the holding time was extended from 60 min to 120 min, no change in the phase composition was detected. The phase composition of the NbC/Fe-based composite layer was relatively stable, and no new phases were formed with the prolonging of the holding time. Figure 3 depicts the cross-sectional SEM images of the samples with different holding times at 1000°C−40 MPa. As can be observed, the niobium plate and HT300 are well combined, with uniform layer thickness, smooth interface, absence of cracks and other type of defects, showing the existence of a good metallurgical bonding. The thickness of the NbC layer was gradually increased with the prolongation of the holding time, and the layer thicknesses were 5 μm,15 μm and 23 μm after the implementation of a holding for 60 min, 90 min and 120 min, respectively.
The EDS spectrum analysis results of the corrosion samples in the reaction zone of 1000°C−40 MPa-120 min are shown in figure 4. It can be observed that an obvious continuous wave-shaped NbC/Fe interface was formed at the interface. This is because the melting point of niobium is 2468°C, whereas the melting point of HT300 is 1270°C. According to the Taiman temperature formula [25] T Taiman = 0.3 ∼ 0.4T M , niobium and iron begin to diffuse at 740°C and 381°C, respectively.
When the temperature was kept at 1000°C, the HT300 and the niobium plate remain solid. The C atoms in the matrix diffused into the niobium plate due to the filling reaction, and as a result, carbide is generated. In the Fe-C-Nb ternary system, the radius ratio of C atoms to Nb atoms was R C /R Nb = 0.91/2.09 = 0.438. According to the Hagg principle [26], C has the ability to diffuse into Nb to form a simple interstitial mutually. Nb is a BCC structure, and the interstitial radii in the 〈100〉 direction of the tetrahedral and octahedral interstitials in the lattice are 0.291R Nb and 0.154R Nb , respectively. These radii are quite small to accommodate C atoms. On the contrary, the interstitial radius in the 〈110〉 direction of the octahedral interstitial is 0.633 RNb, which is enough to accommodate C atoms. Therefore, C atoms are filled in such octahedral interstitials.  By considering the concentration gradient of C from the matrix to the niobium plate, and the fact that the binding force between Nb atoms and C atoms is far greater than the binding force between Fe and C atoms, C atoms diffuse rapidly from the matrix to the metal plate. More metal atoms are relatively dissolved at the interface between the metal plate and the substrate, while C atoms are the most abundant. Moreover, the nucleation rate of carbides is much greater than the growth rate, which leads to the formation of submicron dense ceramics. With the prolongation of the holding time, C atoms are diffused rapidly along the matrix to the metal plate. Hence, the thickness of the NbC/Fe layer is increased.
The interface 1 between the unreacted niobium plate and niobium carbide, and the interface 2 between the gray cast iron and niobium carbide are shown in figure 3(c), where obvious differences can be detected. Interface 2 is well-formed and it is clearly visible. As can be ascertained from figure 4, an obvious interface is formed between niobium carbide and gray cast iron, in which iron and niobium do not diffuse into each other. On the contrary, C atoms diffuse from Fe to Nb through the interface and react with Nb, and a continuous wave-like interface is formed at interface 1.

Microhardness
The microhardness of the samples at 1000°C−40 MPa-90 min and 1000°C−40 MPa-120 min are 2068.4 HV 0.1 and 2096.4 HV 0.1 , respectively, whereas the test results of multiple groups of different holding times are close, indicating that the NbC/Fe composite layer has a uniform structure and its hardness range is comparable to that of pure niobium carbide. (2040 HV ∼ 2550 HV) [27]. Figure 5 displays the hardness comparison of unreacted niobium plate, NbC/Fe-based composite layer and HT300 matrix. As can be observed, the hardness of the NbC/Fe-based composite layer obtained on the surface of HT300 is the highest, and the hardness of the composite layer is 7 times that of the matrix.    figure 6. From the tensile curves, it can be seen that during the process of tensile fracture, with the progress of the loading process, no a necking phenomenon is appeared, which could lead to brittle fracture. At the same time, with the extension of the holding time, a value of 241 MPa was extracted for the tensile strength of the 'sandwich' structure samples at 1000°C−40 MPa-90 min, while and the tensile strength of the samples processed at 1000°C −40 MPa-120 min was up to the value of 297 MPa. In order to further explore the specific fracture position and the interlayer bonding strength, SEM and EDS analyses were performed on both sides of the fracture of a single sample, as is shown in figure 7(c). The morphology and distribution of the niobium carbide ceramic particles can be clearly seen, which are embedded in the iron matrix. In addition, the fracture mode is an intergranular fracture.
The schematic diagram of the tensile fracture is shown in figure 7(a). The energy spectrum analysis of the fractures on both sides of the 1000°C−40 MPa-120 min tensile sample is shown in table 1. Fe (wt%) only accounts for 3.58% and 2.78% of the fractures on both sides indicating that the fracture position is located in the layer of the NbC/Fe composite layer. Thus, it can be argued that the degree of bonding between the composite layer and the iron matrix is greater than the bonding strength inside the composite layer. From the data analysis,   it can be also seen that the bonding strength of the NbC/Fe-based composite layer and the substrate is σ combined 297 MPa. Figure 8 shows the acoustic emission curve and microscopic appearance of scratches at 1000°C−40 MPa-90 min and 120 min. As can be observed, when the critical loads of the NbC/Fe-based composite layer were obtained by holding the heat for 90 min and 120 min are 41.2 N and 75.75 N, respectively, the curve begins to appear noise signal, which indicates the damage of the composite layer. When the load is increased to about 70.05 N and 81.2 N, the intensity of the noise signal is further increased. Although the cracks expand and various types of cracks appear, the overall scratch morphology of the composite layer is complete, and no large-scale collapse effects take place. The microscopic morphology 1 and 2 at the maximum load are shown in figure 9(a), (b).

Microscopic bond strength of NbC/Fe-based composite layer and substrate
From figure 8 it can be argued that the scratch morphology of the samples at 1000°C−40 MPa-90 min and 120 min shows that the scratching process can be divided into three different stages: 0 ∼ Lc 1 mild deformation stage, Lc 1 ∼ Lc 2 moderate deformation failure stage and Lc 2 ∼ 100N heavy damage stage. In the stage of 0 ∼ Lc1, the initial scratch marks are light, the bottom of the groove is smooth, and the surface of the reinforcement layer remains almost unchanged, which is the elastic deformation zone.
In the stage Lc 1 ∼ Lc 2 , a large number of small micro-cracks gradually appeared at the bottom of the scratch furrow and on the edges on both sides. With the increase of the applied load, the number of cracks is increased and the expansion is accelerated. This effect is consistent with the trend of stress change during the scratching  process, which is commonly referred to as the plastic deformation zone. The specific performance is as follows: as the load is increased linearly and the indenter is moved forward, the indenter gradually penetrates the surface of the sample. As a result, a tensile micro-crack perpendicular to the scratch direction is formed behind the area where the indenter is passed through the scratch. At the same time, transverse cracks along the scratch direction began to appear on both sides of the scratch furrow, which were parallel to the scratch direction (figures 9, 4).
In the Lc 2 ∼ 100N stage, when Lc2 is exceeded, with the increase of the load and the penetration depth of the indenter, the transverse annular crack is gradually expanded and the density increases with the loading. At this time, since the energy during the loading process cannot be completely released, the formation of open cracks at an angle of 45°to the scratching direction begins to appear on both sides of the transverse annular crack (figures 9, 3). This effect takes place until the parallel cracks on both sides of the scratch that are perpendicular to the scratching direction extend to the surface of the composite layer outside the scratch (figures 9, 4), and finally, at the edge of the scratch, the composite layer particles peeled off and failed (figures 9, 2).

Simulation of bonding properties within the composite layer
To further study the bonding properties of the NbC/Fe-based composite layer to the HT300 matrix, tensile simulations were carried out under the same conditions as the experiments, as is shown in figure 7(a). According to the load-carrying characteristics of the 'sandwich' structure sample, the sample can be simplified into the matrix part and the NbC/Fe-based composite layer part, and the uniaxial tensile simulation of the composite layer part was carried out. For computational efficiency without loss of generality, the gauge length portion was treated as an equivalent homogeneous material (EHM) [28]. On top of that, to achieve quasi-static conditions, the simulated stretching speed was set to the value of 0.45 mm min −1 .
The simulation of Mises stress and the equivalent plastic strain of the composite layer during the tensile process is shown in figure 10: (a), (b) and (c) present the equivalent stress distribution, while (d) (e) and (f) illustrate the equivalent strain distribution. The distribution of both stress and strain is not uniform. Due to the confinement of the matrix deformation by the reinforcing particles, the deformation of the matrix mainly takes place between the clustered particles, resulting in a 45°increase in both stress and strain along the loading direction, especially near the sharp corners of the particles. During the stretching process, the hard and brittle particles cannot be coordinately deformed with the matrix, causing the development of stress concentration near the sharp corners of the particles, and resulting in the failure of the matrix and the formation of microcracks. Subsequently, the stress and strain values in the matrix rapidly are increased near the crack tip, leading to the failure of the matrix (figure 10(a), (d)). As the applied load is increased, the crack expands along the direction of the maximum shear stress, which also becomes the direction of the stress concentration. During the crack propagation process, the cracks are deflected by the particles. During the process of crack deflection and expansion, the crack always follows the direction of the stress concentration. Finally, the specimen fractures as it propagates until the crack eventually penetrates the entire sample.
As is shown in figure 11, during the deformation and fracture of the selected NbC/Fe composite layer, the particle distribution has a significant impact on the initiation and evolution of cracks. Cracks are always initiated at the particle concentration area or at the sharp corners of the particles. The origins of this effect are associated with the fact that many interfaces exist in the particle concentration area and as a result, the local energy is higher. Therefore, crack initiation is easy to occur during the stretching process. After that, the deformation of the matrix is mainly concentrated at the tip of the particle and the joint between the particle and the matrix. When the equivalent plastic strain reaches a critical value, the failure of the matrix will be induced. The matrix failure propagation is illustrated as microcracks caused by the particle fracture and along the direction of maximum shear stress in the matrix, which is 45°to the loading direction.

Conclusion
In this work, high-purity niobium plate and HT300 were used as raw materials, and the NbC/Fe-based composite reinforcement layer on the surface of the iron matrix was prepared by using the hot pressing diffusion method. The phase composition, microstructure and bonding strength of both the composite layer and the matrix were also systematically studied. The following conclusions can be drawn: (1) At 1000°C−40 MPa, the gray cast iron matrix and the niobium plate have a hot pressing diffusion-reaction, and the NbC/Fe-based composite layer was successfully prepared on the surface of the gray cast iron. With the increase in the holding time, the thickness of the NbC layer became bigger. The layer thicknesses at 60 min, 90 min, and 12 min were 5 μm,15 μm and 23 μm, respectively, the reaction interface was pure without impurities, and no new phases except NbC and α-Fe were formed.
(2) From the scratch test it was demonstrated that the composite layer has a good protection effect on the substrate, whereas no large area failure was detected. The application of a relatively longer holding time leads to better bonding performance of the NbC/Fe-based composite layer. Additionally, the performance against external single abrasive wear was significantly enhanced.
(3) All fracture events of the tensile test occurred in the prepared reinforcing layer. The NbC and the HT300 matrix were well bonded, and the bonding strength between the composite layer and the matrix was greater than the bonding strength within the composite layer, σ combined 297 MPa. From the bonding performance simulated outcomes, it can be argued that the fracture position of the composite layer is located within the composite layer. Interestingly, the appearance of the stress concentration point at the sharp corner of the NbC particles led to the expansion of the crack, indicating that the composite layer is well bonded.