Introduction

Owing to its high hardness, stability at high temperatures, chemical inertness, corrosion resistance, excellent thermal conductivity, and electrical and optical properties, TiN has become a desirable material for tremendous advanced applications in different areas (Ref 1,2,3). The nanocrystalline powder form of TiN can be used as wear-resistance surface protective coating for cutting and machinery tools and for solar control films, and it is also a potential candidate for fusion reactors (Ref 4).

Different routes can be used to prepare TiN such as a sol-gel approach (Ref 5), direct reaction of TiCl4 with ammonia (Ref 6), activated reactive evaporation (Ref 7), a self-propagating combustion method under high nitrogen pressure and high temperature (Ref 8). Chemical vapor deposition (CVD) can also be used to prepare TiN by reacting the TiCl4 with ammonia at temperatures above 1000 °C (Ref 2). In addition, the low-temperature physical vapor deposition (PVD) approach allows the deposition of TiN on metallized wafers (Ref 2). The high costs of these preparations, utilization of toxic precursor gases (e.g., TiCl4, HN3), the toxic exhaust gases (e.g., HCl, Cl), and the excess of contamination in the end-product are considered to be the disadvantages of these methods.

In contrast to the traditional chemical and physical routes used for the preparation of TiN, Calka et al. (Ref 9) and El-Eskandarany et al. (Ref 10) prepared TiN powders at ambient temperature under a flow of nitrogen gas atmosphere using a reactive ball milling (RBM) technique (Ref 11). Since then, a wide variety of metal nitrides and hydrides have been successfully prepared by the RBM technique (Ref 12).

Whereas most of the reported work described the preparation of a few grams (~5 or less) of nanocrystalline TiN powders, the present study proposes a roller-milling technique for the preparation of a large amount (500 g) of nanocrystalline TiN powder upon RBM of polycrystalline of Ti powders under 10 bar of nitrogen gas. The short RBM time required to prepare single-phase fcc-TiN powder with uniform nanocrystalline grains together with the consolidation of the bulk TiN nanophase to full density to achieve high mechanical properties are the advantages of the present work.

Experimental Procedure

Materials

Commercial titanium (Ti) powders with particles 45 μm in diameter and 99.95 wt.% purity, provided by Sigma-Aldrich-GF14536537, USA, as well as, helium (He) and nitrogen (N) gases with purity of 99.999 wt.%, provided by a local gas company in Kuwait were used as starting materials.

Equipment, Devices and Tools

The reactive ball mill vial supplied by ZOZ GmbH, Germany, was made of Cr steel alloy with an internal diameter of 12 cm and capacity of 1200 mL (Fig. 1a and b). The milling media were 4 kg of Cr steel balls 14 mm in diameter and a density of 10 g/cm3 (Fig. 1a). High-pressure gauges and valves were used as evacuation/gas charging kits (Fig. 1c). A horizontal-type roller mill, models RM200-ZOZ GmbH, Germany, was used to carry out the RBM process at room temperature (Fig. 1d). Weighting and handling of the powders were done inside a He gas atmosphere-glove box (UNILAB Pro Glove Box Workstation), provided by mBRAUN, Germany.

Fig. 1
figure 1

Setup of a laboratory-scale roller mill (1200 mL in volume) showing (a) the milling tools including the balls (milling media) and vial, (b) evacuation setup showing gas valve for (c) evacuation and inert gas filling, and (d) ball milling process, using the roller mill. The powder charging kits used in SPS consisted of (e) graphite die and punches, cylindrical graphite spacer and graphite disk spacer. The die charged with the powders was wrapped by carbon felt and tied with carbon yarn (f). A pyrometer used to measure the die’s surface temperature during the SPS process is shown in (g)

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Syntheses of TiN Nanocrystalline Powders

To minimize the Fe contamination introduced to the powders upon ball milling using the FeCr milling tools, 5 g of Ti was first milled with the FeCr balls under an Ar gas atmosphere for 50 h. After performing this milling run, the milling tools (the vial and the balls) were completely coated by thin Ti layers due to the cold welding process that occurred during the milling process. 50 g was massed inside the glove box and then sealed together with 400 Cr steel balls (Fig. 1a) using a ball-to-powder weight ratio of 80:1. The sealed vial was then removed from the glove box (Fig. 1b) and evacuated to 10−3 bar before pressurizing under 10 bar of nitrogen gas (Fig. 1c). The RBM process started by mounting the nitrogen-pressurized vial onto a roller mill (Fig. 1d) operated at ambient temperature with a rotation speed of 350 rpm. The processing RBM time varied from 1 to 20 h. The RBM process was periodically halted after selected milling times (1, 3, 6, 12, and 20) where 2 g of the ball-milled powder was discharged from the vial inside the glove box. The vial was then pressurized again with nitrogen gas before resuming the RBM process. This RBM process was repeated 6 times to ensure the reproducibility of the end-product for each stage of milling. However, the powder of the end-product obtained after 20 h of RBM time was collected from 5 individual runs, giving a total amount of approximately 223 g.

Powder Consolidation

The spark plasma sintering (SPS) system, provided by Dr. Sinter Lab. Instrument, Japan, was employed to consolidate the powder samples obtained after each stage of the RBM runs (1, 3, 6, 12, and 20 h). A schematic of the SPS system used in the present study is illustrated in Fig. 2. The consolidation system consists of a sintering press unit with vertical single-axis pressurization, specially designed punch electrodes incorporating a water cooler, a water-cooled vacuum chamber, a vacuum/air/argon-gas atmosphere control mechanism, a special direct current (DC) pulse sintering power generator, a cooling-water control unit, a Z-axis position measuring and control unit, temperature measuring and control units, an applied pressure display unit, and various safety interlock devices.

Fig. 2
figure 2

A schematic illustration of SPS used in the present study. The system consists of an SPS sintering machine with a vertical single-axis pressurization mechanism, specially designed punch electrodes incorporating a water cooler, a water-cooled vacuum chamber, a vacuum/air/argon-gas atmosphere control mechanism, a special DC-pulse sintering powder generator, a cooling-water control unit a position measuring unit, a temperature measuring unit, an applied pressure display unit and various interlock safety units

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In the typical SPS process, approximately 5 g was charged into a 15 mm graphite die and stacked between the upper and lower punches (Fig. 1e). To ensure unchallenging ejection of the sample after sintering and in order to avoid any unexpected reactions between the internal surfaces of the die tools (die and punches) and the sample, graphite sheets were used as spacers (Fig. 1e). On the other hand, the radiant heat transfer to the machine was minimize by enclosing both the entire die and punch assembly with carbon felt, using carbon yarn as shown in Fig. 1(f). The die was then mounted on the sintering stage in the SPS chamber and held between the upper and lower punch electrodes (Fig. 2).

All powder samples without exceptions were consolidated in vacuum at 1600 °C under a vertical pressures ranging between 19.6 to 38.2 MPa for 5 min. The total amount of the powder samples obtained after each RBM time was consolidated into 6 bulk samples with equal dimension of diameter (15 mm) and aspect ratio (1:1). During the sintering process, the temperature of the die was monitored and measured, using a pyrometer system, as displayed in Fig. 1(g).

It should be emphasized that the actual processing temperature of powder consolidation by SPS can be quite different from the measured value by the pyrometer (Ref 12, 13). This is because the pyrometer measures the temperature at a niche in the die which is neither exactly on the sample surface nor in the surface interior. In addition, the measured temperature presents an average value and gives no indication of the local temperature that can actually exist between the particles.

Characterization Techniques

To monitor the progress of the gas-solid reaction conducted in the RBM vial, all of the ball-milled Ti powder samples obtained after various stages of RBM were carefully analyzed by different techniques. The average crystal structure of all samples was investigated by means of x-ray diffraction (XRD) with CuKα radiation, using a 9 kW Intelligent x-ray diffraction system, provided by SmartLab-Rigaku, Japan. The local structure of the ball-milled powders at the nanoscale was studied by 200 kV field emission high-resolution transmission electron microscopy/scanning transmission electron microscopy (HRTEM/STEM) system supplied by JEOL-2100F, Japan. The local (~5 nm) composition of the nanocomposite powder samples was investigated by energy-dispersive x-ray spectroscopy (EDS, Oxford Instruments, UK) equipped with the HRTEM/STEM system. For the bulk samples obtained by SPS, TEM samples were prepared by the ION Slicer, EM-09100 IS system, provided by JEOL to get a circular sample with a diameter of 3 mm and a thickness of 100 nm. The morphological properties of the powders obtained after different ball milling times and their corresponding SPSed bulk samples were investigated by a 15 kV field emission scanning electron microscope (FE-SEM, JSM-7800F, Japan). The hardness of the compacted samples (~10 mm in diameter) was investigated, using a conventional Vickers indenter with a load of 5 kg. The hardness values reported subsequently are averaged from 60 indentation measurements obtained from 6 individual samples for each RBM time. The wear properties indexed by the coefficient of friction for the as-consolidated bulk samples were investigated at ambient temperature in dry mode, using a pin-on-disk machine, provided by Microtest, Spain). A 10-mm-diameter ball made of composite WC/Ni was used as pin, where the applied load and disk speed were kept constant at 200 N and 150 rpm, respectively. The densities of the consolidated samples were determined by Archimedes’ principle, using water immersion. Some mechanical properties of the consolidated samples were determined by nondestructive testing using an impulse excitation technique, ASTM-E-1876-99 (IMCE nv, Belgium).

The chemical analyses of the powder samples before and after consolidation were performed in order to determine the concentrations of N, O, and the Fe contamination contents. The Fe content was analyzed by inductively coupled plasma optical emission spectrometry (ICP), where the amount N and O were determined with gas chromatograph (GC) system, using a plasma emission detector.

Results and Discussion

Structural Changes with Changing RBM Time

XRD pattern of the starting elemental hcp-Ti powder samples milled under 10 bar of nitrogen gas atmosphere for different RBM time is presented in Fig. 3(a), (b), (c), and (d). The powder obtained after 1 h consisted of large crystallites, indicated by the sharply narrowed Bragg peaks corresponding to hcp-Ti, as identified by the Miller indices of the (100), (002), (101), and −(110) shown in Fig. 3(a). After this early stage of RBM time, a small mole fraction of the TiN phase was detected, as indicated by the very low intensity Bragg peaks related to fcc-TiN (#PDF file 01-071-9845), as presented in Fig. 3(a). The powder samples at this early stage of milling experienced lattice imperfections, exemplified by the nanotwins, and stuck faults in their internal structure, as indexed by the circular symbols shown in Fig. 3(e). The HRTEM micrograph taken at the near edge of an aggregated powder particle obtained after 3 h of RBM revealed the existence of reacted TiN phase embedded into a Ti matrix, as elucidated in Fig. 3(f).

Fig. 3
figure 3

XRD diffraction patterns of Ti powder samples obtained after RBM for (a) 1 h, (b) 6 h, (c) 12 h, and (d) 20 h. The HRTEM images taken for the samples milled for 1 h and 3 h are shown in (e) and (f), respectively

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During the intermediate stage of RBM (6 to 12 h), the Bragg peaks corresponding to the TiN phase became pronounced (Fig. 3b and c), indicating an increase in the mole fraction of the reacted TiN phase. The HRTEM micrograph of the powder obtained after 6 h of RBM time shows two TiN particles, with diameter of approximately 20 nm on the unprocessed Ti matrix zone (Fig. 4a). Both the particles and matrix suffered from significant amounts of lattice imperfections implied by the presence of stacking faults and dislocations, as denoted by the circular symbols shown in Fig. 4(a). As the RBM time increased (12 h), the intensity of the Bragg peaks of unreacted hcp-Ti decreased, implying a decrease in the mole fraction of Ti metal, as presented in Fig. 3(c).

Fig. 4
figure 4

HRTEM images of the samples obtained after (a) 6 h and (b) 20 h. The FFT of the selected circle shown in (b) is displayed as an inset of (b). The STEM-DFI of the powders obtained after 20 h of RBM is displayed in (c) where the corresponding EDS map of elemental Ti and N is shown in (d) and (e), respectively

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After 20 h of RBM, all the Bragg peaks corresponding to unprocessed metallic Ti powders disappeared, a single phase of nanocrystalline TiN appeared, characterized by broad Bragg peaks displayed in Fig. 3(d). We should emphasize that those Bragg peaks corresponding to the fcc-TiO2 phase came from the local oxidation of TiN powder surfaces during the preparations of XRD samples. The lattice parameter ao of the TiN phase obtained after 20 h (end-product) was calculated from (111) and (200) reflections and found to be 0.423 nm, matching well with the reported value of pure TiN (0.4235 nm, PDF# 01-071-9845).

The HRTEM image of TiN powder samples obtained after 20 h of RBM is presented together with fast Fourier transformations (FFT) in Fig. 4(b). Obviously, the powder consisted of fine grains of about 8 nm in diameter (or less) with different crystallographic orientation, as shown in Fig. 4(b). Moreover, the corresponding FFT of the selected zone denoted by the ring symbol (Fig. 4b) implied the existence of TiN grains oriented to the zone axis [200]. Under this resolution of FE-HRTEM, no hcp-Ti crystals could be detected in the powder of the end-product, suggesting the completion of the gas-solid reaction via the RBM process.

Local Compositional Analysis

The homogeneity of the TiN powders obtained after 20 h of RBM time was intensively examined by EDS, using the HRTEM/STEM technique. The STEM-dark field (DF) image of the near edge of an aggregated TiN particles is shown in Fig. 4(c). Based on the corresponding EDS elemental mapping of Ti (Fig. 4d) and N (Fig. 4e), it can be concluded that both elements were homogeneously distributed in the examined area (~100 nm × 150 nm) without elemental segregation beyond the nanoscale level, as presented in Fig. 4(d) (Ti-Ka1) and 4e (N-Ka1). Moreover, the average composition (in at.%) of this end-product calculated from EDS spot analysis of 36 different zones was Ti51.68N48.32.

Morphology

Selected FE-SEM micrographs of the powder samples obtained after the early, intermediate, and final stages of RBM are shown together in Fig. 5. The sample obtained after 1 h of milling was bulky in shape and had wide particle size distribution, ranging between 20 to 110 μm, as displayed in Fig. 5(a). Increasing the RBM time to 6 h led to a considerable decrease in the particle size and the formation of significant large volume fractions of particles with spherical morphology, as presented in Fig. 5(b). The local-spot EDS analyses of those large particles in Fig. 5(b) show negligible values of N content; however, the spherical fine particles (under 10 μm) possessed nearly equiatomic compositions of TiN. Increasing the RBM time to 12 h led to increasing the fine TiN particles (less than 1 μm in diameter) against the unprocessed Ti (~1 to 1.7 μm in diameter) powders, as shown in Fig. 5(c). Toward the completion of RBM processing time (20 h), the end-product TiN powders possessed very smooth surfaces, and spherical morphology with narrow particle size distribution in the range between 190 to 400 nm in diameter, as shown in Fig. 5(d).

Fig. 5
figure 5

FE-SEM micrograph of the powder samples obtained after (a) 1 h, (b) 6 h, (c) 12 h, and (d) 20 h of RBM

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Chemical Analysis

Chemical analysis was performed in order to determine the N, O, and Fe contents in the ball-milled powder samples obtained after different stages of RBM. As the RBM time increased, the N content monotonically increased to reach approximately 18 wt.% after 12 h of RBM (Fig. 6). The N absorbed by the Ti powder reached to the level of 19.7 wt.% after 20 h of RBM and saturated at this value even after longer RBM time (30 h), as shown in Fig. 6. This implies the completion of the gas-solid reaction via the RBM process and the formation of TiN. The O content detected in the powders is independently changed with changing the RBM time, as shown in Fig. 6. The O contamination was attributed to handling the powder during preparations of the samples for chemical analysis outside the glove box. The end-product obtained after 20 h of RBM was contaminated with approximately 1.4 wt.% O, as shown in Fig. 6.

Fig. 6
figure 6

Nitrogen, oxygen, and Fe contents on the milled powder samples obtained after different stages of RBM time

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As mentioned in the experimental section, the milling tools were coated with Ti layers in order to minimize the Fe introduced to the powders upon the RBM process. Despite this, a small amount of Fe content (0.09 to 0.85 wt.%) was introduced to the Ti powders during the early stage of milling (3 to 6 h), as shown in Fig. 6. Increasing the RBM time to 12-20 h led to increasing the volume fraction of TiN (hard phase), resulting in an obvious increase in the Fe contamination from the balls due to the abrasion, as presented in Fig. 6. After 20 h, the maximum Fe contamination in the TiN powders was approximately 2 wt.%, as elucidated in Fig. 6.

Consolidation with SPS Technique

The bright field image (BFI) of ion-sliced SPS-consolidated powder samples obtained after 1, 12, and 20 h of RBM are shown in Fig. 7. The bulk sample of the powder milled for 1 h (Fig. 7a) consisted of large grains (~300 nm in diameter) of polycrystalline hcp-Ti, as confirmed by the corresponding selected area diffraction pattern (SADP) shown in Fig. 7(b). As previously shown (Fig. 3c), the powders obtained after 12 h of RBM were composed of unreacted Ti and fully reacted TiN fine powders. The bulk consolidated sample corresponding to these powders revealed a wide grain size distribution (50 to 280 nm in diameter), as shown in Fig. 7(c). The SADP taken from the center of Fig. 7(b) confirms the presence of Ti and TiN phases, as shown in Fig. 7(d). Consolidation of the end-product TiN powders obtained after 20 h of RBM into a bulk sample did not lead to severe grain growth, as characterized by the fine nanocrystalline grains ranging between 42 to 86 nm in diameter shown in Fig. 7(f). Comparing the grain size of the consolidated bulk TiN compact with the original size obtained after 20 h of RBM (8 nm (Fig. 4b)), we can conclude that the SPS can only lead to a moderate growth of TiN grains by approximately 5 times. Moreover, the SPS consolidation step did not led to the formation of any other reactive phase rather than the original fcc-TiN phase, as confirmed by the SADP displayed in Fig. 7(g).

Fig. 7
figure 7

BFI and corresponding SADP of the planar view for as-consolidated and then ion-sliced powder samples obtained after (a, b) 1 h, (c, d) 12 h, and (f, g) 20 h of RBM

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In fact, the powder sintering principles in the SPS process are based on the presence of an electric field called the field-assisted sintering technique (FAST). In contrast with the conventional sintering techniques in which the sample is heated from the outside, the sintering procedure in SPS takes place upon heating the sample internally by the passage of an electric current. When SPS is compared with the conventional hot pressing technique, one can realize that the SPS process is taking place with extremely high heating rates. Moreover, external pressures, which can also be applied during the sintering process in the range between few MPa up to 1000 MPa, lead to a rapid and effective consolidation process. SPS is an extraordinary rapid sintering process in which the whole process (including temperature rise and holding times) are completed with a short period ranging from 5 to 20 min. This direct way of heating allows the application of very high heating and cooling rates, enhancing densification over grain growth promoting diffusion mechanisms, maintaining the intrinsic properties of nanopowders in their fully dense products.

Although, intensive research efforts have been carried out regarding the consolidation of metal, ceramics, and their composites by SPS, the mechanism of this sintering procedure is still not well established. As mentioned in the above section, the main difference between SPS and other sintering approaches is that both die and powder are directly heated by the Joule effect of the DC current. The ON-OFF DC-pulse energizing method generates spark plasma, spark impact pressure, joule heating, and an electrical field diffusion effect (Ref 14). In this process, the surfaces of the powders are activated and purified when compared with the other traditional sintering methods. Fresh powder surfaces lead to an easy sintering process, characterized by a lower sintering temperature and a shorter processing time. Upon starting the ON-OFF DC current path, the pulse current flows through the powder particles that are loaded into conductive graphite sintering tools (die and punches). The pulsed current has a cleaning effect on the surfaces of the powder particles based on the observation of grain boundaries without oxidation formed between particles (Ref 15). Such an intense joule heating effect at the surface of the conductive powder particles results in reaching the boiling point and therefore leads to the localized vaporization or cleaning of powder surfaces (Ref 16). Moreover, it has been suggested by Tiwari et al. (Ref 16) that joule heating resulting from the passage of electric current through the powder particles assists and enhances the welding of the particles under mechanical pressure.

Physical and Mechanical Properties

Density

Figure 8(a) presents the bulk densities taken from 6 consolidated independent samples obtained after different milling times. The density of the samples obtained after 1 to 3 h of RBM varied from 4.1 to 4.8 g/cm3, as shown in Fig. 8(a). This wide variation was attributed to the composition variations of the powders after this early stage of milling containing large volume fractions of pure Ti against a small fraction of the processed TiN reacted phase. As the RBM time increased (6 to 12 h), the density of the consolidated sample tended to have larger value (4.5 to 4.9 g/cm3), suggested an increasing in the volume fraction of the dense TiN phase (5.4 g/cm3) against the light Ti phase (4.5 g/cm3), as shown in Fig. 8(a). After 20 h of RBM, the density ranged between 5.29 to 5.37 g/cm3 with an average value of 5.34 g/cm3. Comparing this value with the theoretical reported value of pure TiN (5.38 g/cm3, PDF# 01-071-9845), we can conclude that SPS of RBMed powders can led to the formation of a nearly full dense TiN bulk material, with a relative density of 99.26%. The FE-SEM micrograph of the fracture surface for the powders consolidated after 20 h of RBM shows a very dense-structure morphology, implying the formation of a high-density fully developed TiN bulk material, as displayed in Fig. 8(b). Increasing the RBM time to 30 h did not show a significant variation on the density value that remained in the range of 5.27 to 5.35 g/cm3, as shown in Fig. 8(a).

Fig. 8
figure 8

Effect of RBM time on (a) the bulk density and (c) Vickers hardness for the powders obtained after different stages of milling. The FE-SEM micrograph of the fracture surface and Vickers indentation for the consolidated samples obtained after 20 h of RBM are shown in (b) and (d), respectively

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Vickers Hardness

Vickers hardness measurements taken from 6 consolidated independent samples obtained after different milling time are displayed in Fig. 8(c). The hardness of those samples obtained after completion the early stage of RBM (1 to 3 h) varied from 1.9 to 7.2 GPa, as shown in Fig. 8(c). Increasing the RBM time led to increasing the volume fraction of the hard TiN phase against the unreacted Ti metal. This can be understood by looking at the hardness values, which were varied from 4 to 20.2 GPa, as shown in Fig. 8(c). Toward the end of the RBM time (20 to 30 h), the hardness values of the consolidate samples were almost saturated at an average value of 22.9 GPa ±1.25, as shown in Fig. 8(c). This implies the formation of single TiN phase and the disappearance of Ti metal. The FE-SEM micrograph of the Vickers indentation for the SPSed sample obtained after 20 h of RBM is elucidated in Fig. 8(d). The micrograph reveals obvious cracks emanating from the indentation corners, implying brittle behavior of the obtained TiN phase.

Elastic Modulus

The elastic modulus values, indexed by the Young’s and shear moduli of the consolidated samples obtained after different RBM times, were estimated from the measured sample densities and the constant parameters of the nondestructive testing apparatus. Figure 9 shows the effect of RBM time on the Young’s and shear moduli. These values increased significantly during the early and intermediate stages of RBM time (1 to 12 h), suggesting an increase of TiN phase in the milled powder. Toward the end of the RBM time (20 to 30 h), both Young’s and shear values are almost saturated at 384 and 192 GPa, respectively, indicating the formation of a single homogeneous TiN phase.

Fig. 9
figure 9

Influence of RBM time of the Young’s and shear modulus values of the powder samples milled for different stages of milling

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Wear Properties

The relation between the coefficient of friction (μ) and the sliding distance of as-consolidated powder samples obtained after different RBM times is presented in Fig. 10. The wear test was achieved using pin-on-disk technique, using a WC-Ni composite ball (pin) under the application of 200 N, and a rotation speed of 150 rpm. All samples had similar behaviors of reaching steady-state stages after approximately 20 m of sliding distance, as shown in Fig. 10. During the second stage (20 to 100 m), the μ value for 1 h sample was dramatically increased to 0.65 after 100 h, as displayed in Fig. 10. The samples obtained after 6 h and 12 h of RBM that contain larger TiN volume fractions showed slightly increased values after 100 m of approximately 0.49 and 0.41, respectively (Fig. 10). The sample obtained after 20 h of RBM (end-product) showed a lower μ value of approximately 0.35, as displayed in Fig. 10. It can be then concluded that, in addition to the attractive high hardness value demonstrated for the nanocrystalline TiN bulk material, this system also possesses superior wear-resistance behavior, making it a desirable material for a surface protective coating employed for cutting tool applications in oil and underground water sectors.

Fig. 10
figure 10

Dependence of coefficient of friction (μ) on the RBM time and sliding distance

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Conclusion

Based on the results of the present study, a single reacted phase of TiN nanopowders was synthesized by the RBM approach, starting from pure Ti powder samples and nitrogen gas atmosphere. A roller mill operated at 350 rpm was operated at room temperature under 10 bar of a nitrogen gas atmosphere. The obtained TiN powders after 20 h of RBM possessed spherical morphology with an average diameter of 0.5 μm and were composed of nanocrystalline grains with an average diameter of 8 nm. To investigate some of the physical and mechanical properties of the powders obtained after different stages of RBM time, the samples were consolidated under vacuum at 1600 °C into cylindrical bulk compacts 15 mm diameter, using the SPS technique. The relative density of the end-product was above 99% of the theoretical density. Moreover, this consolidation step did not lead to drastic grain growth and the TiN bulk material maintained its nanocrystalline characteristics with an average grain size of 56 nm in diameter. The synthetic nanocrystalline TiN bulk material obtained after 20 h of RBM revealed a high Vickers hardness of the end-product for TiN, which was measured and found to be 22.9 GPa. Moreover, the measured Young’s and shear modulus values of this end-product bulk TiN were 384 and 189 GPa, respectively. The nanocrystalline TiN bulk material shows also excellent wear resistance with a coefficient of friction of 0.35.