Journal of Materials Research and Technology Journal of Materials Research and Technology
Original Article
Influence of Mo alloying on the thermal stability and hardness of ultrafine-grained Ni processed by high-pressure torsion
Garima Kapoora, Yi Huangb, V. Subramanya Sarmac, Terence G. Langdonb, Jenő Gubiczaa,,
a Department of Materials Physics, Eötvös Loránd University, P.O.B. 32, Budapest H-1518, Hungary
b Materials Research Group, Faculty of Engineering and the Environment, University of Southampton, Southampton SO17 1BJ, UK
c Department of Metallurgical and Materials Engineering, Indian Institute of Technology Madras, Chennai 600036, India
Received 11 April 2017, Accepted 25 May 2017

The influence of Mo alloying on the thermal stability of grain size, dislocation density and hardness of ultrafine-grained (UFG) Ni alloys was studied. The UFG microstructure in alloys with low (∼0.3at.%) and high (∼5at.%) Mo contents was achieved by high-pressure torsion (HPT) performed for 20 turns at room temperature. The thermal stability of the two alloys was studied by calorimetry. A Curie-transition from ferromagnetic to paramagnetic state was not found for the Ni–5% Mo alloy due to the high Mo content. It was found that heating at a rate of 40K/min up to ∼850K resulted in a complete recovery and recrystallization of the UFG microstructure in the alloy with 0.3% Mo. The same annealing for Ni–5% Mo led only to a moderate reduction of the dislocation density and the grain size remained in the UFG regime. Therefore, the higher Mo content yielded a much better thermal stability of the Ni alloy. The influence of the change of the microstructure during annealing on the hardness is discussed.

High-pressure torsion, Ni–Mo alloys, Dislocations, Grain size, Hardness, Thermal stability
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Ni–Mo alloys possess high hardness and wear resistance which make them useful for various practical applications. For instance, these alloys can be used as hard coating materials [1]. In addition, they exhibit corrosion resistance to the reducing acids, such as hydrogen chloride. Thus, Ni–Mo alloys can be used as catalysts in the production of hydrogen as these acids induce reduction reactions and generally result in hydrogen evolution at cathodic sites [2–4]. These alloys also show high activity and long-term stability as hydrogen evolution reaction catalysts under alkaline conditions. They are also applicable as substrate material for superconducting coatings which can be produced by severe cold rolling and subsequent annealing [5]. The latter processing provides extremely sharp cube texture as required for epitaxial coatings.

Improvement in the mechanical properties of Ni–Mo alloys can be attained by the implementation of efficient and novel severe plastic deformation processes, such as equal-channel angular pressing (ECAP), cryorolling or high-pressure torsion (HPT). Over the years, there is an increasing interest in employing SPD techniques for production of ultrafine-grained (UFG) materials [6–11]. HPT is considered as the most effective SPD method in grain refinement and improvement of the strength of metallic materials [12–17]. The enhanced properties of UFG microstructures can be degraded due to grain growth at elevated temperatures [18–22]. The stability of UFG microstructures is an important aspect for their reliable performance in practical applications. Thus, a study of microstructure stability of UFG materials by annealing to high temperatures is crucial.

The addition of alloying elements further influences the development of the UFG microstructures during SPD as well as their thermal stability. Solute atoms have pinning effects on lattice defects, such as dislocations and grain boundaries, therefore they can stabilize the UFG microstructure [23–28]. The influence of concentrations of alloying elements on the thermal stability of face-centered cubic (fcc) metals (e.g., Ni and Cu) processed by SPD has been investigated in the literature [29,30]. However, despite the significant possible applications of Ni–Mo alloys, no investigations have been conducted to date to examine the influence of annealing on the microstructure and properties of HPT-deformed Ni–Mo alloys. Therefore, the present study was initiated to study the influence of heat-treatment on the microstructure and mechanical properties of UFG Ni alloys with two different Mo concentrations. The samples processed by a two-step combination of cryorolling and 20 turns of HPT were annealed to about 850K and the microstructural parameters were compared with the HPT-processed state. In addition, the variation of the hardness during annealing was examined and correlated to the change of microstructure.

2Experimental material and procedures2.1Processing of UFG Ni–Mo alloys

Two Ni alloys with low and high Mo contents were processed by induction melting and casting into a Cu-mold. The chemical compositions of the as-processed Ni alloys were determined by energy dispersive spectroscopy (EDS) in a scanning electron microscope (SEM). The alloys with ∼0.28 and ∼5.04at.% of Mo were labeled as low-Mo and high-Mo, respectively. Although, in addition to Mo, other elements, such as Al (0.84–1.08at.%), Fe (0.13–0.25at.%) and Si (0.05–0.34at.%) were also found in the samples, the major difference between the chemical compositions of the two samples was the much higher Mo content in the material labeled as high-Mo. Detailed chemical compositions for low-Mo and high-Mo alloys were given earlier [31]. The as-cast ingots with diameters of ∼32mm were hot-rolled at 1100°C to a thickness of ∼13mm. The hot-rolled samples were then subjected to a combination of cryorolling and HPT. First, small strips cut from the hot-rolled materials were processed by cryorolling at liquid nitrogen temperature (LNT). This procedure resulted in a reduction of the thickness from ∼13mm to ∼3mm in multiple passes with a reduction of ∼5% per pass. Then, disks with diameters of ∼10mm and thicknesses of ∼1mm were prepared from the cryorolled materials. In the second step of SPD, the samples were processed by the HPT technique under quasi-constrained conditions [32] with an applied pressure of 6.0GPa and a rotating speed of 1rpm at RT for 20 turns.

2.2Differential scanning calorimetry

Our former study [31] revealed that for both alloys the microstructural parameters and the hardness were saturated between the half-radius and the periphery of the disks processed by 20 turns of HPT. Therefore, the thermal stability of the microstructures was investigated in the region between the half radius and periphery of the disks using differential scanning calorimetry (DSC). The DSC experiments were conducted on small samples cut from these regions of the HPT disks by a Perkin Elmer (DSC2) calorimeter at a heating rate of 40K/min under an Ar atmosphere. Then, the HPT-processed samples were heated up to a characteristic temperature of the DSC thermograms. These specimens are regarded as annealed samples.

2.3Microstructure from EBSD

The microstructures of the HPT-processed and the annealed Ni alloys were studied by electron backscatter diffraction (EBSD) using an FEI Quanta 3D SEM. Before EBSD, the specimens were mechanically polished first by SiC abrasive papers with 600, 1200, 2500 and 4000 grit and then by a colloidal silica suspension (OP-S) having a particle size of 1 micrometer. After mechanical polishing, the surfaces were treated by electro-polishing at 28V and 1A using an electrolyte with a composition of 70% ethanol, 20% glycerine and 10% perchloric acid (in vol.%). The step size in the EBSD study was ∼30nm and the grain sizes were evaluated using Orientation Imaging Microscopy (OIM) software. It is noted that only those regions in the EBSD images which were bounded by high-angle grain boundaries (HAGBs) with misorientations higher than 15° were considered as grains.

The distortions inside the grains were analyzed using Kernel Average Misorientation (KAM) maps prepared by the OIM software. In this evaluation process, a local misorientation angle value was assigned to each pixel, which was determined as the average misorientation between the studied central pixel and all pixels at the perimeter of the kernel around the investigated pixel. The radii of the kernels were ∼50nm for all images even if their pixel sizes were smaller in order to make their KAM maps comparable.

2.4Characterization of microstructure by X-ray diffraction

The phase composition and the microstructural parameters were investigated by X-ray diffraction (XRD). The Debye–Scherrer diffraction rings were detected on two-dimensional imaging plates using a high-resolution diffractometer (Multimax-9 made by Rigaku) with CuKα1 radiation (wavelength: λ=0.15406nm). The X-ray line profile analysis (XLPA) was carried out with the Convolutional Multiple Whole Profile (CMWP) fitting method in order to determine the dislocation density [33]. In this fitting procedure, the diffraction pattern is fitted by the sum of a background spline and the convolution of the instrumental pattern and the theoretical line profiles related to dislocations, crystallite size and planar faults. As an example, Fig. 1 shows the CMWP fitting for the high-Mo sample processed by 20 turns of HPT. The open circles and solid line represent the measured and fitted patterns, respectively.

Fig. 1.

CMWP fitting of the X-ray diffraction pattern taken at the half radius of the high-Mo alloy disk processed by 20 turns of HPT. The open circles and the solid line represent the measured data and the fitted curve, respectively. The intensity is in a logarithmic scale.

2.5Microhardness testing

The Vickers microhardness of the HPT-processed and the annealed samples was measured using a Zwick Roell ZHμ hardness tester with an applied load of 0.5kg and a dwell time of 10s for both compositions of Ni alloys.

3Experimental results

Illustrative EBSD images for the HPT-processed microstructures of both alloys were presented in our previous study [34]. Figs. 2 and 3 show the grain sizes and the dislocation densities for low-Mo and high-Mo samples obtained immediately after 20 turns of HPT. The grain size of the HPT-processed high-Mo alloy was smaller (∼130nm) than the value obtained for the low-Mo sample (∼184nm). In addition, the dislocation density was about two times higher (∼60×1014m−2) for the former material than for the latter specimen (∼23×1014m−2). The higher dislocation density and the smaller grain size for the HPT-processed high-Mo alloy can be attributed to the hindering effect of Mo atoms on the annihilation of dislocations and the motion of grain boundaries during deformation.

Fig. 2.

Histogram of the grain sizes in the samples processed by 20 turns of HPT and the specimens annealed at ∼850K for both low-Mo and high-Mo alloys.

Fig. 3.

Histogram of the dislocation densities in the samples processed by 20 turns of HPT and the specimens annealed at ∼850K for both low-Mo and high-Mo alloys.

Fig. 4a and b shows the DSC thermograms obtained at 40K/min for the HPT-processed low-Mo and high-Mo samples, respectively. For both materials, the first and the second heating scans are also shown. For the low-Mo material, an endothermic peak was observed for both scans with a peak temperature of ∼627K. This peak is related to the Curie-transition from a ferromagnetic to paramagnetic state. Between ∼600 and ∼830K, the heat flow for the first heating scan is considerably larger than for the second scan, indicating exothermic processes. The maximum of this endothermic peak can be found at ∼743K. This peak corresponds to the recovery and/or recrystallization of the UFG microstructure in the HPT-processed low-Mo alloy, as shown in the next paragraph. For the high-Mo material, the endothermic peak related to the ferromagnetic to paramagnetic transition was not observed. However, a long exothermic peak was detected between ∼630 and ∼1000K. The peak maximum can be found at about 786K. The characteristic temperatures of the DSC thermograms are listed in Table 1.

Fig. 4.

DSC thermograms obtained at a heating rate of 40K/min for the regions between the half radius and the periphery of the disks processed by 20 turns of HPT for (a) low-Mo and (b) high-Mo alloys.

Table 1.

The characteristic temperatures of the thermograms obtained for the low- and high-Mo alloys by DSC scan at 40K/min. The Curie-temperature of the high-Mo alloy was taken from Ref. [35].

  Low-Mo  High-Mo 
Curie-temperature [K]  627  333 [35] 
Temperature range of the exothermic DSC peak [K]  600–830  630–1000 
Exothermic peak maximum [K]  743  786 

In order to characterize the difference between the thermal stability of the HPT-processed low-Mo and high-Mo alloys, additional samples were heated up to ∼850K at 40K/min and then quenched to RT. This temperature was selected as above this temperature there is no exothermic signal for the low-Mo alloy (see Fig. 4a). The grain size and the dislocation density of the annealed low-Mo and high-Mo samples were determined by EBSD and XLPA, respectively. Fig. 5a and b shows inverse pole figure (IPF) maps obtained by EBSD for the annealed low-Mo and high-Mo alloys, respectively. The HAGBs and the LAGBs are indicated by black and white lines, respectively. The grain size of the annealed samples determined from the IPF maps are shown in Fig. 2. For the low-Mo material, the average grain size increased from ∼184 to ∼874nm while in the case of the high-Mo sample the grain size was enhanced from ∼130 only to ∼225nm. In addition, for the low-Mo sample most of the boundaries are straight and there are no orientation gradients in the grains which indicate that the lattice distortion is negligible. The KAM maps in Fig. 5c and d show the local misorientations between 0 and 5° for the HPT-processed low-Mo and high-Mo alloys after annealing at ∼850K. The KAM maps are indicative of local strains in the studied microstructures. It can be seen that for the annealed low-Mo alloy considerable local misorientations may be observed only in the vicinity of some grain boundaries.

Fig. 5.

IPF maps obtained by EBSD for the HPT-processed (a) low-Mo and (b) high-Mo alloys after annealing at ∼850K. The color code for the maps is shown in the inset in (a). The HAGBs and the LAGBs are indicated by black and white lines, respectively. KAM maps showing the local misorientations between 0 and 5° for the HPT-processed (c) low-Mo and (d) high-Mo alloys after annealing at ∼850K.

These results suggest that the microstructure in the low-Mo alloy was recrystallized during annealing and this is supported by the practically zero dislocation density. In fact, XLPA evaluation could not be performed for the annealed low-Mo sample as the peak breadth was as narrow as the instrumental profile. Thus, the dislocation density in this material was lower than the detection limit for the applied diffractometer configuration (∼1013m−2). The Debye–Scherrer ring for reflection 220 of the annealed low-Mo specimen is shown in Fig. 6a. The spotty ring also suggests that the sample exhibits a recrystallized microstructure. Each high intensity spot is scattered from a recrystallized grain. Fig. 6b shows that the Debye–Scherrer ring for the annealed high-Mo alloy is homogeneous, which indicates that this sample was not recrystallized at ∼850K. The dislocation density in high-Mo decreased from ∼60×1014m−2 to ∼19×1014m−2 so that only a moderate recovery occurred during annealing. This is confirmed by the KAM map in Fig. 5d, which reveals high lattice distortion inside the grains even after annealing at ∼850K. It is noted that in the low-Mo alloy the recrystallization also resulted in an elongated shape for many grains (see Fig. 5a) since the nucleated new grains are preferably separated from the parent grains by low energy boundaries growing only on specific crystallographic planes (e.g., twin boundaries on planes {111}). At the same time, in the annealed high-Mo alloy the grain shape remained equiaxed due to the lack of recrystallization (see Fig. 5b).

Fig. 6.

Debye–Scherrer rings for reflection 220 obtained for the HPT-processed (a) low-Mo and (b) high-Mo alloys after annealing at ∼850K.

The higher thermal stability of the high-Mo alloy is also reflected in the variation of hardness during annealing (see Fig. 7). While the hardness decreased considerably from ∼3300 to ∼1340MPa during annealing of the HPT-processed low-Mo alloy, a significant change in the very high hardness of ∼4320MPa for the high-Mo specimen was not observed. The reduction of the hardness for low-Mo can be explained by the considerable decrease in the dislocation density and the grain-growth that occurred due to recrystallization (see Figs. 2 and 3). For the high-Mo sample, the much lower decrease of the dislocation density and the moderate increase of the grain size suggests a lower reduction of the hardness. Therefore, the practically unchanged hardness for the high-Mo alloy is surprising and seems to be in contradiction with the changes of the microstructure. However, this effect can be explained by the plastic deformation occurring during the hardness measurement, as discussed in the next section.

Fig. 7.

Histogram of the hardness values for the samples processed by 20 turns of HPT and the specimens annealed at ∼850K for both low-Mo and high-Mo alloys.


Fig. 4 shows that while the Curie temperature for low-Mo was detected as ∼627K, in the case of the high-Mo alloy an endothermic signal for a ferromagnetic to paramagnetic, transition was not observed. This observation is in line with former studies which indicated that the Curie temperature of Ni significantly decreased with increasing Mo content [35]. The literature data in [35] show that for ∼5at.% (8wt.%) Mo alloying there was a reduction of Curie temperature to 333K, therefore, the paramagnetic transition cannot be observed in the present experiments starting at 300K, due to the transient DSC signal in the beginning of the measurements. The Curie temperature for low-Mo (∼627K) agrees with the value characteristic for pure Ni, so that the low-Mo content of 0.3at.% has no considerable effect on the ferromagnetic–paramagnetic transition.

Our study demonstrates that the thermal stability of the present HPT-processed UFG Ni alloys is strongly influenced by the Mo content. While the low-Mo material fully recrystallized at ∼850K, in the high-Mo alloy only a partial recovery and a moderate grain growth were observed. With increasing Mo concentration, the exothermic DSC peak related to recovery and recrystallization was shifted to a higher temperature. Moreover, while this peak ends at ∼850K for the low-Mo sample, the heat release finishes only at ∼1000K for the high-Mo alloy. The much better thermal stability of the HPT-processed UFG microstructure for the high-Mo material can be attributed to the retarding effect of the Mo solute atoms on the motion of dislocations and grain boundaries, which are necessary for recovery and recrystallization. In addition, former studies (e.g., [36]) have shown that the segregation of alloying elements on grain boundaries may reduce the energy of boundaries, thereby reducing the driving force for recrystallization and grain growth. As a consequence, the segregation of solute atoms can stabilize UFG microstructures.

Although the higher Mo content resulted in a much better thermal stability, the dislocation density decreased to about half of the value achieved after HPT and the grain size also increased by about 50% during annealing. Therefore, it was a surprising result that the hardness remained practically unchanged despite the recovery and grain growth. This phenomenon can be understood if we consider the influence of the hardness measurement on the probed material. Namely, during hardness testing a probe tip is pressed into the material which results in a plastic deformation around the indent. During this plastic deformation, dislocations are generated so that the hardness value characterizes this strain hardened material. The equivalent plastic strain corresponding to the hardness measurement with a Vickers or Berkovich indenter is about 8% [34]. Therefore, the hardness value is related to the flow stress at a plastic strain of 8% and not to the yield strength of the material before hardness testing. For severely deformed metallic materials, the strain hardening is usually low and therefore the difference between the yield strength and the flow stress at a plastic strain of 8% is very small. Thus, for these specimens the hardness value can be related to the yield strength. At the same time, for annealed samples the strain hardening is usually large and, therefore, the difference between the yield strength and the flow stress at a plastic strain of 8% is significant. As a consequence, in the latter cases the hardness cannot be related directly to the annealed microstructure. In the present study, for the high-Mo sample the annealing must have resulted in a decrease of the yield strength due to recovery and grain growth. However, the flow stress at a plastic strain of 8% may be considerably larger than the yield strength so that the hardness cannot reflect the change in the microstructure during annealing. In addition, besides the density of dislocations, their arrangement has also an influence on the dislocation strengthening effect. Specifically, a more clustered dislocation structure is usually associated with a larger strengthening effect for the same dislocation density [37,38]. XLPA evaluation also determines parameter M, which reflects the arrangement of the dislocations. A smaller value of M relates to a more shielded strain field of the dislocations, and the arrangement of dislocations into low energy configurations, such as LAGBs or dipoles, yields a consequent decrease in M. The present experiments show that for the HPT-processed high-Mo alloy the value of M decreased from ∼4.8±0.5 to ∼1.6±0.2 during annealing up to ∼850K. The more clustered dislocation structure after annealing decreases the softening effect of the reduced dislocation density. Most probably, this effect and the strain hardening caused by the hardness testing resulted in the unchanged hardness after annealing.

5Summary and conclusions

The influence of Mo content on the thermal stability of UFG microstructures in Ni alloys processed by HPT was studied. The following conclusions were drawn:

  • 1.

    The exothermic peak observed in the DSC thermograms was shifted to a higher temperature with increasing Mo content. For a low-Mo alloy, this peak finished at ∼830K while for a high-Mo material the exothermic processes, such as recovery and recrystallization, end only at ∼1000K. In addition, alloying with ∼0.3at.% Mo in Ni had only a negligible effect on the Curie temperature, while for ∼5at.% Mo a ferromagnetic–paramagnetic transition was not observed in the studied temperature range.

  • 2.

    Annealing up to ∼850K resulted in a full recrystallization of the UFG microstructure in the low-Mo alloy. In this process, the grain size increased from ∼184 to ∼874nm while the dislocation density decreased from ∼23×1014m−2 below 1013m−2. At the same time, for the high-Mo material only a partial recovery and moderate grain growth was observed. In this material, the dislocation density decreased from ∼60×1014m−2 to ∼19×1014m−2 and the grain size increased from ∼130 to ∼225nm. The better thermal stability of the high-Mo alloy can be attributed to the pinning effect of Mo solute atoms on dislocations and grain boundaries which retards their motion during recovery and recrystallization. In addition, the segregation of Mo at grain boundaries reduces the boundary energy, thereby decreasing the thermodynamical driving force for recrystallization and grain-growth.

  • 3.

    The hardness for the low-Mo alloy considerably decreased from ∼3300 to ∼1340MPa due to recrystallization during annealing. At the same time, the partial recovery and the moderate grain growth failed to yield significant hardness reduction for the high-Mo material. This phenomenon can be explained partly by the clustering of dislocations into a denser structure as revealed by the reduced dislocation arrangement parameter. In addition, the hardness measurement causes a plastic deformation around the indenter which results in strain hardening in the annealed sample.

Conflicts of interest

The authors declare no conflicts of interest.


This research was presented at the 3rd Pan American Materials Congress held as part of the TMS Annual Meeting in San Diego, California, on February 27–March 2, 2017. The authors are grateful to Mr. Alajos Ö. Kovács and Mr. Gábor Varga for DSC and EBSD investigations, respectively. This research was supported by the Hungarian Scientific Research Fund, OTKA, Grant no. K-109021. Two of the authors (YH and TGL) were supported by the European Research Council under ERC Grant Agreement No. 267464-SPDMETALS.

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Paper was a contribution part of the 3rd Pan American Materials Congress, February 26th to March 2nd, 2017.

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