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Johnson Matthey Technol. Rev., 2021, 65, (4), 584

doi:10.1595/205651321x16123456338021

Microstructure Evolution of Ruthenium During Vacuum Hot Pressing

Obtaining fine-grained metal with mechanical properties close to electron-beam melted ruthenium

  • Renyao Zhang, Junmei Guo, Chuanjun Wang, Limin Zhou, Ming Wen*
  • State Key Laboratory of Advanced Technologies for Comprehensive Utilization of Platinum Metals, Kunming Institute of Precious Metals, Kunming 650106, China
  • *Email: wen@ipm.com.cn
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Article Synopsis

Ruthenium tablets with mean grain size of ~4–5 μm were prepared by vacuum hot pressing (VHP), and tablets with maximum density of 12.2 g cm–3 were obtained with sintering time of 2 h. X-ray diffraction (XRD) revealed that there was a texture change with sintering time. The microstructure of the ruthenium tablets was observed by electron backscatter diffraction (EBSD) and field emission scanning electron microscopy (FSEM). The microstructure evolution of ruthenium with sintering time is discussed.

1. Introduction

The refractory platinum group metal ruthenium exhibits unique properties such as high melting point (2334°C) (1), conductivity (1.348 × 107 Ω m) (2) and high hardness (~337 DPN for as-melted ruthenium surface) (3). Ruthenium has been used as an active catalyst in applications such as ammonia synthesis and chemical water splitting (4, 5). Ruthenium-based thin films have gained considerable research interest especially in the electronics industry. They have been widely used as electrode materials for dynamic random access memory (DRAM) (6), perpendicularly magnetised heterostructures (7) and as a seed layer material for copper interconnects or transparent conductive zinc oxide (810). These uses can be attributed to the low resistivity of ruthenium, its relatively high work function and its low reactivity with various metals.

Magnetron sputtering via a ruthenium sputtering target is a well-known technique for ruthenium film deposition since the deposition process provides excellent productivity and is widely used for mass production (1114). Ruthenium sputtering targets having a homogeneous fine-grained structure are vital for the preparation of high-quality ruthenium films. The powder metallurgy (PM) technology technique of VHP is used for manufacturing ruthenium sputtering targets. To the best of the present authors’ knowledge, although some work has been done on the residual stress of ruthenium sintered by spark-plasma-sintering (SPS) and VHP (15, 16), there are very few published works on the structural evolution of ruthenium during VHP which is quite important for the industrial application of ruthenium sputtering targets. In the present work, the structure and microhardness of ruthenium tablets prepared by VHP are examined and discussed.

2. Material and Methods

2.1 Sample Preparation

High-purity ruthenium powder (99.995 wt%) with an average particle size of 5 μm provided by the Kunming Institute of Precious Metals, China, was used as the raw material. Ruthenium samples were compacted under a pressure of 40 MPa at 1250°C in a vacuum of 10–3 Pa at a heating rate of 15°C min–1 for 0.5 h, 1 h, 2 h, 3 h and 4 h. After sintering, the power was turned off, the tablets were cooled in the furnace to room temperature before being taken out. The prepared samples had the shape of a tablet with a diameter of 30 mm and a height of 4 mm.

2.2 Microstructure and Properties Characterisation

The density of the samples was measured by Archimedes’ method (17). The phase content of the samples was examined with the help of the X’Pert PRO X-ray diffractometer (PANalytical, The Netherlands) and the SmartLab 9 Kw (Rigaku Corporation, Japan) operated by copper Kα irradiation. The morphology of the ruthenium samples and their fracture surfaces were studied on the Sirion 200 field emission scanning electron microscope and the Versa 3DTM DualBeamTM (FEI Company, USA) including EBSD. The fracture surface was prepared as follows: firstly, a ruthenium tablet of 10 × 4 × 0.5 mm was cut by electric discharge machining (EDM); then the disc was polished by abrasive paper until the thickness was ~200 μm; finally, the thin slice was bent by hand. A Leica EM TIC 3X ion beam milling system (Leica Microsystems GmbH, Germany) was used to prepare the sample’s surface for EBSD analysis. The Vickers microhardness of the samples was measured with the help of the HXS‐1000A microhardness tester (Shanghai Highwell Optoelectronics Technology Co, Ltd, China) with a load of 100 g.

3. Results and Discussion

3.1 Microstructure Characterisation

The density results are shown in Figure 1. It can be seen that the sample density increases with sintering time between ~0.5–2 h, then decreases after 4 h. The maximum density of ruthenium was 12.2 g cm–3 with a sintering time of 2 h. The XRD spectra of the ruthenium samples are given in Figure 2. Ruthenium powder exhibits random orientation of the particles corresponding to the standard PDF card of ruthenium (PDF #06-0663) (18). The intensities of and peaks decrease sharply, the intensity of (0002) peak increases, while the intensities of (0002) and peaks are the same in the samples sintered for 0.5 h. The samples sintered for 1 h show similar XRD patterns. This indicates that a stable grain structure forms in the ruthenium samples after 1 h. The intensity of peak began to increase after sintering for 2 h and 4 h. The intensity of peak increased again in the sample sintered for 4 h. A stable (0002) texture was formed in the ruthenium samples sintered for 0.5–2 h; however, it disappeared after 4 h. This finding has shown that the texture is controlled by varying the sintering time.

Fig. 1.

The density results of ruthenium samples

The density results of ruthenium samples

Fig. 2.

The XRD spectra of the sintered ruthenium samples (Cu Kα irradiation)

The XRD spectra of the sintered ruthenium samples (Cu Kα irradiation)

Ruthenium powder is shown in Figure 3. It consists of particles having irregular shapes with sizes from 1 μm to 15 μm (Figure 3(a)). The large particles are aggregates of the small particles. There are many pores or voids in the particles (Figure 3(b)), which were caused by gas released during the chemical reduction of ruthenium. The apparent density and the tap density of ruthenium powder were 1.9 g cm–3 and 3.2 g cm–3, respectively.

Fig. 3.

The morphology of the ruthenium powder at: (a) low magnification and (b) high magnification

The morphology of the ruthenium powder at: (a) low magnification and (b) high magnification

Morphology of the fracture surfaces of ruthenium samples are given in Figure 4. Their fracture mode is attested as brittle intergranular fracture (19). This finding agrees with the conclusion that ruthenium behaves like a brittle solid even at elevated temperatures (20). This is a puzzling behaviour because a hexagonal close-packed (HCP) metal is a ductile material (21). The slip of dislocations on both prismatic and basal planes could happen in ruthenium single crystals at room temperature (22). However, basal slip is the main deformation mode of ruthenium single crystals at room temperature under tension and their fracture mode is brittle transgranular fracture (23). The grains in the samples are homogeneous having a size of ~4–5 μm, which does not depend on the sintered time.

Fig. 4.

The fracture surface of ruthenium samples: (a) sintered for 0.5 h, low magnification; (b) sintered for 0.5 h, high magnification; (c) sintered for 1 h, low magnification; (d) sintered for 1 h, high magnification; (e) sintered for 2 h, low magnification; (f) sintered for 2 h, high magnification; (g) sintered for 4 h, low magnification; (h) sintered for 4 h, high magnification

The fracture surface of ruthenium samples: (a) sintered for 0.5 h, low magnification; (b) sintered for 0.5 h, high magnification; (c) sintered for 1 h, low magnification; (d) sintered for 1 h, high magnification; (e) sintered for 2 h, low magnification; (f) sintered for 2 h, high magnification; (g) sintered for 4 h, low magnification; (h) sintered for 4 h, high magnification

In order to further reveal the sintering mechanism, EBSD was used to analyse the microstructure. Figure 5 is the grain boundary map of ruthenium. Firstly, it can be seen that the grain size of the ruthenium samples with different sintering time is ~4–5 μm, corresponding well with the fracture surface (Figure 4). The grain interior shows different morphologies with sintering time. For ruthenium sintered for 0.5 h, there are few twins and low-angle grain boundaries (~5–15°) in the grain interior (Figure 5(a)). For ruthenium sintered for 1 h, Figure 5(b) shows clean grain structure with fewer low-angle grain boundaries as compared with Figure 5(a). Twins and low-angle grain boundaries with high density appear inside grains when sintering time reaches 2 h (Figure 5(c)), and then the defect density decreases again in the sample sintered for 4 h. The statistical results of grain boundaries are summarised in Table I. From Table I, it can be seen that the total length of low-angle grain boundaries is only 1.8 μm for ruthenium sintered for 1 h which is the shortest of all the samples. The total length of low-angle grain boundaries is 42 μm for ruthenium sintered for 2 h, which is the longest of all the samples. The total length of boundaries increases from ~0.5–2 h, achieving a maximum value with sintering time of 2 h (263 μm), and then decreases to 141 μm with sintering time of 4 h. For some HCP metals, the activation of twins shows a strong dependence on the deformation temperature and strain rate. During the hot compression of titanium at temperatures from 673 K to 973 K, Zeng et al. found many twins at 723 K and 0.1 s–1, few twins after deformation at 723 K and 0.01 s–1 and no twins after deformation at 973 K and 0.01 s–1 (24).

Ruthenium possesses an HCP lattice, where twinning is the important stress accommodation channel under mechanical loading (21). Experiment has shown that there are four active twin systems: , , and in the ruthenium samples. Figure 6 is the twin boundary map of ruthenium sintered for different times at 1250°C. For ruthenium sintered for 0.5 h, it can be seen there are a few twin systems in the grain interior (Figure 6(a)). Figure 6(b) shows that for ruthenium sintered for 1 h, there are fewer twins in the grain interior, and the percentage of twin systems decreases while the other twin systems of , and increase as compared with Figure 6(a). The total twin density increases again, and the percentage of twin systems increases when sintering time reaches 2 h (Figure 6(c)). The total twin density decreases again with a sintering time of 4 h, although the percentage of twin systems increases to a maximum value of 72.1%. The statistics of twin boundaries in ruthenium samples are summarised in Table II. It can be seen that the longest twin was 42.6 μm for ruthenium sintered for 2 h. The total length of twin boundaries increases from ~0.5–2 h while the total length of twin boundaries achieves a maximum value of 64.3 μm with a sintering time of 2 h, then the value decreases to 44.2 μm with a sintering time of 4 h. In HCP titanium, the ratio of the lattice constants (c:a) is 1.587, which is similar to that of ruthenium (c:a = 1.582) (23). Previous research has shown that deformation twins occur during hot compression of titanium at 723 K and 0.1 s–1 (24).

Fig. 5.

The grain boundary map of ruthenium sintered at different times at 1250°C: (a) sintered for 0.5 h; (b) sintered for 1 h; (c) sintered for 2 h; (d) sintered for 4 h. Red line: 5°<angle<15°, green line: 15°<angle<30°, blue line: 30°<angle<100°

The grain boundary map of ruthenium sintered at different times at 1250°C: (a) sintered for 0.5 h; (b) sintered for 1 h; (c) sintered for 2 h; (d) sintered for 4 h. Red line: 5°<angle<15°, green line: 15°<angle<30°, blue line: 30°<angle<100°

Fig. 6.

The twin boundary map of ruthenium sintered at different times at 1250°C: (a) sintered for 0.5 h; (b) sintered for 1 h; (c) sintered for 2 h; (d) sintered for 4 h. Red line: The twin boundary map of ruthenium sintered at different times at 1250°C: (a) sintered for 0.5 h; (b) sintered for 1 h; (c) sintered for 2 h; (d) sintered for 4 h. Red line:  twins, green line:  twins, purple line:  twins, blue line:  twins twins, green line: twins, purple line: twins, blue line: twins

Table I

Statistics of the Grain Boundary Types in Ruthenium Samples

Time, h 5°–15°
15°–30°
30°–100°
Length, μm Percentage, % Length, μm Percentage, % Length, μm Percentage, %
0.5 5.6 5.5 14.4 14.4 80.2 80.1
1 1.8 1.6 11.2 9.9 100.4 88.5
2 42.0 16.0 58.0 22.0 163.1 62.0
4 25.2 17.9 14.0 9.9 101.8 72.2
Table II

Statistics of Twin Boundary Types in Ruthenium Samples

Time, h



Length, μm Percentage, % Length, μm Percentage, % Length, μm Percentage, % Length, μm Percentage, %
0.5 14.4 51.0 6.0 21.3 3.6 12.6 4.3 15.1
1 10.3 30.7 9.3 27.6 6.6 19.7 7.4 22
2 42.6 66.3 7.2 11.2 4.5 7.0 9.9 15.5
4 31.8 72.1 4.9 11.2 3.7 8.3 3.7 8.4

3.2 Hardness Characterisation

To examine the effect of microstructure on the mechanical properties of ruthenium, the Vickers microhardness of the samples was measured. The dependence of the hardness of the ruthenium samples on the sintering time is shown in Figure 7. The hardness increases at first and then it decreases with sintering time. The hardness of ruthenium sintered for 0.5 h was 447.2 HV, and it increased to a maximum hardness of 540.1 HV for ruthenium sintered for 1 h. The hardness decreased to 531.6 HV for ruthenium sintered for 2 h, then decreased to the minimum value of 407.8 HV for ruthenium sintered for 4 h. In a previous study of ruthenium hardness (3) the sintered tablets exhibited a hardness between 91 HV to 377 HV, while after hot working their hardness became 307 HV to 455 HV. The difference in the measurements may be explained by the fact that the samples in (3) were sintered without pressure and, as a result, their density was lower (9.69–11.88 g cm–3).

Fig. 7.

The hardness of the ruthenium samples with sintering time

The hardness of the ruthenium samples with sintering time

3.3 Discussion

It was shown that the average grain size in the ruthenium samples is stable (~4–5 μm) and does not depend on the time of sintering at the process temperature of 1250°C (Figure 4). It seems that sponge particles could recrystallise under annealing at 1250°C for ~0.5–4 h, while this temperature is sufficiently low that the grains in the samples could begin growing during this short time. The initial powder size of materials may impact the grain size of the tablets during high pressure-high temperature sintering. For example, Shin et al. found that during the sintering of diamond there was no abnormal grain growth (AGG) for the initial powder size of 4 μm, while AGG happened for the initial powder size of 2 μm (25). Thus, the particle size of the present initial ruthenium powder may be suitable for the present sintering conditions.

Early in the sintering process, after a sintering time of 0.5 h (Figure 5(a)), the pressure and high temperature caused particle rearrangement, sliding and metallurgical bonding. There were a few twins in some grain interiors suggesting the initial inhomogeneous deformation of ruthenium. The inhomogeneous state may be caused by different grain orientations in which some orientations deform easily or by areas with closely spaced grains and metallurgical bonding which deform first. As the sintering time reached 1 h (Figure 5(b)), the grains showed clean and uniform grain structure. In VHP samples, the particles were pressed together and kept in contact (26). Thus, there were more diffusion paths to promote atomic migration and induce sintering in multiple directions (27). With the appropriate pressure, temperature and holding time, voids were further crushed and ruthenium particles contacted and bonded with each other to form fully metallurgical bonding across grain boundaries. The density of ruthenium also increased slightly from 0.5 h to 1 h.

With the increase of sintering time to 2 h, ruthenium formed metallurgical bonding in almost the whole bulk material (Figure 5(c)). Defects such as twins and low-angle grain boundaries appearing inside grains point to some plasticity in the ruthenium samples. The total length of grain boundaries and twin boundaries reached a maximum value for a sintering time of 2 h, implying the maximum plastic deformation for the present ruthenium sample. Thus, the density also reached a maximum value after sintering for 2 h. As sintering time further increased to 4 h (Figure 5(d)), there were fewer defects and boundaries inside the grains as compared with ruthenium sintered for 2 h, and the density also reached a minimum value in all samples sintered between ~0.5–4 h.

It is well known that annealing can induce crystal formation from such defects as dislocations and twins. Hence, the decreasing hardness of the samples sintered for 4 h could be caused by the annealing of twins which appeared in the material at earlier stages of sintering. During VHP, the ruthenium powders sustained the crushing of voids, formation of grain boundaries, grain growth, plastic deformation (formation of defects) and recovery of defects. As for the hardness of ruthenium with sintering time (Figure 6), the hardness increased first from 0.5 h to 1 h due to the crushing of voids and formation of tight boundaries. It was found that the maximum hardness was achieved for ruthenium sintered for 1 h, while the density reached its maximum value for ruthenium sintered for 2 h. According to Figure 1 and Figure 7, the samples’ density begins decreasing when sintered for 4 h. The hardness is prone to similar behaviour after processing for more than 0.5 h. This is normal behaviour for hardness because the twins’ density decreases after 4 h. It can be seen from the XRD pattern of ruthenium that the strongest peak is (0002) plane for ruthenium sintered for 1 h, while the strongest peak is plane for ruthenium sintered for 2 h (Figure 1). Since there was little difference in density and average grain size for the ruthenium samples sintered for 1 h and 2 h, the crystal orientation and defect density of ruthenium may play a role in the determination of hardness.

Ruthenium is a brittle metal due to its anisotropic HCP crystal structure that provides a limited number of independent slip systems, and twinning is an important deformation mode. Twinning has been observed in some metals and ceramics such as tungsten carbide, cubic boron nitrides and aluminium oxynitride ceramic during high-pressure and high-temperature sintering (2830). Previous research has shown that deformation of ruthenium occurs by slipping on and twinning on , , and occasionally (3). In the present work, slipping and twinning occurred during VHP of ruthenium even at 1250°C, and the main twin system was . The transition of the main twin system to may be attributed to the high temperature and pressure. Any HCP structure is anisotropic in comparison with a cubic structure. As a result, the total plasticity of an HCP-metallic single-crystal depends on its crystallographic orientation. It is significant for zinc and cadmium at room temperature because only the basal slip is active in these metals under these conditions. The orientation anisotropy of plasticity in titanium and zirconium is not so visible, insomuch as the prismatic slip’s contribution is added to the basal slip. The contribution of twinning to the total plasticity of an HCP-metal is minor, while twinning could influence its hardness and, perhaps, work-hardening. Also, it should be noted that a lot of twin systems exist in HCP-metals, which are known to be ductile and malleable materials. The low malleability of ruthenium is connected with the low cohesive strength of grain boundaries. In the case of PM ruthenium, it is due to brittle intergranular fracture, whose source is likely to be non-metallic impurities rather than its intrinsic properties. Further work still needs to be done to understand the relationship between impurity elements and the plasticity of ruthenium.

4. Conclusions

The VHP technique for the manufacture of ruthenium tablets based on powder metallurgical technology has been described. The method described here allows fine-grained metal with mechanical properties close to electron-beam melted ruthenium to be obtained.

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Acknowledgments

The authors gratefully acknowledge support of the National Key Research and Development Program (No. 2017YFB0305503), International Science and Technology Cooperation Foundation of Yunnan Province, China (No. 2014IA037), and Yunnan Innovation Team Project (No. 2019HC024).

The Authors


Renyao Zhang graduated from Yunnan Agricultural University, China, in June 2015 with a Bachelor’s degree in engineering. He is currently a Master’s student of Kunming Institute of Precious Metals, China, and conducts research under the guidance of Wen Ming. At present, the main research fields are precious metal sputtering targets.


Junmei Guo received her Master’s degree in Materials Science from Kunming Institute of Precious Metals. She has been working for Kunming Institute of Precious Metals since 1994. Her research interest is precious metal alloys and thin films.


Chuanjun Wang received his Master’s degree in Materials Science from Kunming Institute of Precious Metals. He has focused on the research and development of rare and precious metal sputtering targets for more than 11 years.


Limin Zhou graduated from Kunming University of Science and Technology, China, in 2000 with a Bachelor’s degree in engineering. He has worked for Kunming Institute of Precious Metals since 2005. His research interest is new preparation methods of precious metal alloys.


Ming Wen received his PhD degree in Materials Science from the School of Materials Science and Engineering of Shanghai Jiaotong University, China. He has worked for Kunming Institute of Precious Metals since he graduated from Shanghai Jiaotong University. His research interest is sputtering targets and related thin films. He has published more than 20 papers in Science Citation Index (SCI) journals recently.

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