Effect of thermomechanical processing via rotary swaging on properties and residual stress within tungsten heavy alloy

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Effect of thermomechanical processing via rotary swaging on properties and residual stress within tungsten heavy alloy ⁎

Lenka Kunčickáa, Adéla Macháčkováb, Nicholas P. Laveryc, Radim Kocichb, , Jonathan C.T. Cullenc, Libor M. Hlaváčd Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Žižkova 22, 61600 Brno, CZ Faculty of Materials Science and Technology, VŠB-Technical University of Ostrava, 17. listopadu 15, 70833, Ostrava 8, CZ c Materials Research Centre, College of Engineering, Swansea University Bay Campus, Fabian Way, Swansea, Neath Port Talbot, SA1 8QQ, UK d Faculty of Electrical Engineering and Computer Science, VŠB-Technical University of Ostrava, 17. listopadu 15, 70833, Ostrava 8, CZ a

b

A R T I C LE I N FO

A B S T R A C T

Keywords: Tungsten heavy alloys Finite element analysis Stress/strain measurements Electron microscopy X-ray analysis

The effects of cold and warm rotary swaging and subsequent post-process annealing on mechanical properties, residual stress, and structure development within WNiCo powder-based pseudo-alloy were predicted numerically and investigated experimentally. The swaging temperature of 900 °C imparted increase in the Young’s and shear moduli; the post-process annealing at 900 °C also imparted decrease in the residual stress values, primarily due to structure recovery introduced within the matrix. Cold rotary swaging at 20 °C imparted ultimate tensile strength of almost 1 900 MPa, while warm rotary swaging at 900 °C introduced increased plasticity (almost 24 % after a single swaging pass). Post-process heat treatment promoted diffusion of W to the Ni/Co matrix, which increased strength, but remarkably decreased elongation to failure and residual stress. Numerically predicted results of mechanical behaviour corresponded to the experimental results and confirmed the favourable effects of the selected thermomechanical treatments on WNiCo performance.

1. Introduction Given by their high density and the advantageous combination of high strength and sufficient ductility, tungsten heavy alloys (THA) are widely used for demanding applications such as kinetic energy penetrators and radiation shields, or in the space industry. THAs typically contain between 90 and 97 wt. % of tungsten plus other alloying elements, the combination and volume fractions of which significantly influence the strength and plastic properties of the final product. The most common THAs are the W-Ni-Fe [1], W-Ni-Co [2], and W-Ni-Fe-Co [3] systems. Probably the most widespread THA fabrication technology enabling the achievement of enhanced mechanical properties combines methods of powder metallurgy (mechanical alloying, sintering, etc. [4–6]). However, green-sintered tungsten alloys cannot meet high demands due to their poor mechanical (especially tensile) properties. By this reason, plastic deformation processing under hot or cold conditions is usually performed as the subsequent step [7]. Among the forming methods applied to increase the strength and utility properties of THAs are for example cold rolling and its variants [8,9], drawing [10], swaging [11], extrusion [12,13], methods of severe plastic deformation

⁎

(SPD) such as equal channel angular pressing (ECAP) and its modifications and high pressure torsion (HPT) [14–19], and rotary swaging (RS) [20,21]. THAs can also advantageously be fabricated via specialized technologies, such as spark plasma sintering (SPS) [22], and by additive manufacturing, e.g. selective laser melting (SLM) [23]. However, the effects of the preparation procedures on enhancement of the mechanical and physical properties of THAs are not equivalent. Among the key factors influencing the final THA structures and properties are processing parameters (the temperatures applied during sintering, subsequent plastic deformation, and possible heat treatment), but also intrinsic material properties (the ability of the material to form adiabatic shear bands, or tendency to form precipitates), all of which can introduce heterogeneity possibly leading to the development of residual stress, structure instability, and eventual material failure [24]. THAs’ mechanical properties depend on the strength of both the W and γ matrix phases, and on W-W and W- γ interfacial properties, too. The comparison of mechanical properties – ultimate tensile strength and impact toughness – for W-Ni-Co based THAs with different ratios of the individual elements was performed by Kumari et al. [25], who found that microstructures with higher volume fractions of matrix and higher distances between W particles featured superior tensile strength and

Corresponding author. E-mail address: [email protected] (R. Kocich).

https://doi.org/10.1016/j.ijrmhm.2019.105120 Received 13 August 2019; Received in revised form 23 September 2019; Accepted 30 September 2019 0263-4368/ © 2019 Elsevier Ltd. All rights reserved.

Please cite this article as: Lenka Kunčická, et al., International Journal of Refractory Metals & Hard Materials, https://doi.org/10.1016/j.ijrmhm.2019.105120

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impact toughness, regardless the chemical composition. Increasing W content within the matrix enhances the mechanical properties by introducing solid solution strengthening, increasing the matrix volume fraction, and reducing W-W contiguity. Similar effect was observed for the increasing content of Co. Ravi Kiran et al. [26] compared the mechanical properties of THAs with 93 % of tungsten and varying alloying elements ratios (Ni, Fe, Co and Re) and found the alloys featuring high Co contents to exhibit the highest plasticity, while a high Re content resulted in high tensile strength. Another affecting factor is the applied aging temperature, which can significantly alter not only the mechanical properties. As documented by a detailed fractographic study [3], the W-Ni-Fe-Co alloy fracture mode changed from inter-granular to trans-granular after swaging and subsequent aging treatment, depending on the selected thermo-mechanical processing regime: trans-granular failure of W grains was predominant within samples aged at 500 °C, whereas samples aged at 700 °C exhibited ductile tear. THAs are also prone to exhibit Dynamic Strain Ageing/Portevin Le-Chatelier's effect, as proven by Das et al. [27], who documented that swaging with 20–40% deformation at temperatures ranging between 400 and 500 °C was associated with notable serrated yield. Such behaviour can be attributed to the presence of interstitial atoms [28], especially carbon [29]. Nevertheless, studies researching the effects of deformation temperature on the mechanical and physical properties of THAs are scarce. Generally, high processing temperatures can lead to the development of structure instabilities, grain growth, precipitation/recrystallization phenomena, oxidation, and also to degradation of plastic properties of THAs. On the other hand, low forming temperatures can cause technological issues during processing due to the high flow stress of tungsten and its problematic formability resulting from its low number of available slip systems. Therefore, optimization of the applied thermomechanical treatment from the viewpoints of processing temperatures and deformation ratios is of the utmost importance. The primary aim of this study was to investigate the effects of cold and warm rotary swaging supplemented with post-process heat treatment on changes in physical and mechanical properties of the analysed WNiCo pseudo-alloy. Thorough examination via ultrasound measurements, X-ray diffraction analyses, scanning electron microscopy observations, and tensile and microhardness measurements were performed to evaluate the individual swaging routes from the viewpoint of their suitability for production of THA bars. The experimental investigations were supplemented with numerical prediction of mechanical behaviour of the swaged-pieces, including internal stress development.

Fig. 1. Schematic depiction of swaged-piece and rotary swaging dies, sensors and planes evaluated via numerical simulation depicted by arrows.

changes in their structures and properties. The annealed samples swaged via routes A and B were then denoted as A1 and B1 samples, respectively. The thermal and mechanical history, as well as possible porosity, are known to affect the Young’s modulus and can also be the cause of a higher standard error [33,34]. Besides, the material properties of Young’s and shear modulus, Poisson’s Ratio, and density were necessary to perform the XRD texture examination needed to determine the σ11 σ22 residual stress and strain for each sample. By these reasons, determination of changes of the physical properties after the applied thermomechanical processing was among the essential experimental steps. The elasticity properties were measured using an ultrasound method with an average of 5 measurements taken per sample. The ultrasound method assumes that the sample is homogenous and isotropic with errors of 0.1% and 0.5% for longitudinal and shear directions, respectively. The samples for structure development characterization were mechanically ground, polished using vibratory polishing and scanned using a TESCAN Lyra 3 FIB/SEM with a NordlysNano EBSD detector (IPM, CAS). The substructures and grains analyses were performed using ATEX [35], and Channel5 software (scans were taken in ¼ diameter distance from the surface to be comparable to the results of the numerical prediction). Optical images of the samples’ surfaces were acquired via a KEYENCE VHX7000 equipment. TEM images were acquired on ion polished thin foils with a JEOL 2100F device. Possible presence of residual stress was determined via internal grains misorientations analyses performed on EBSD scans and supplemented by X-Ray Diffraction (XRD) measurements. The internal grains misorientations were calculated relatively from the average value and evaluated in the low angle boundaries range, i.e. 0-15°. The XRD was performed using a BRUCKER D8 Advance XRD machine (Swansea University, Wales, UK). For subsequent stress-strain analyses, the samples were measured within the angle range of 127.0001° to 134.9953° with an increment of 0.0298°. The investigated sample was rotated and tilted in order to observe residual stress within the crystal lattice; ϕ was 0°, 45°, 90° and 180° and ψ ranged from 0° to 45° with 5° intervals (Fig. 2). The applied stress model was Biaxial + shear, where σ33 stress tensor is stated to be zero. The strain was calculated by measuring the change of crystal length (Eq. (2)), and the stress in ϕ direction was calculated by (Eq. (3)) [36],

2. Materials and methods 2.1. Experimental methods Preparation of the studied WNiCo (92.6 wt. % W, 5 wt. % Ni and 2.40 wt. % Co) alloy was performed the same way as in our previous study [30]. As the subsequent technological step, the sintered pieces with the diameters of 30 mm were subjected to the technology of rotary swaging [31,32], the schematics of which is depicted in Fig. 1. Swaging was performed via two swaging passes in total at two different temperatures of 20 °C (route A) and 900 °C (route B). The particular temperature of warm swaging was selected in order to minimise possible grain growth and oxidation phenomena. The overall applied deformation degree of 0.81 was calculated using Eq. (1), where S0, Sn are initial and final work-piece cross-sections, respectively.

φ = ln

S0 Sn

εψ =

d ϕψ − do (2)

do

where do is original crystal length under no strain, and dψϕ is crystal length in direction of ϕψ, ϕ is angle between σ1 and σ2, and ψ is angle between σ2 and σ3,

(1)

σϕ =

Swaged-pieces processed via both the routes were subsequently subjected to post-process heat treatment – annealing at 900 °C for 15 minutes in protective Ar2 atmosphere – in order to characterize possible

E ⎛ d ψ − dn ⎞ (1 + ѵ) sin2 ψ ⎝ dn ⎠ ⎜

⎟

(3)

where E is Young’s modulus, ѵ is Poisson’s ratio, dn is crystal length along normal to surface, dψ is crystal length in ψ direction. 2

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Fig. 2. Schematic depiction of positioning the XRD measured sample.

experimental ones. The deformation behaviour was characterized via evaluation of the distribution of effective strain across the longitudinal cutting plane, as well as via evaluation of the development of effective strain during the entire swaging pass in three monitored points (sensors) positioned on the longitudinal cutting plane (as seen in Fig. 1b). Localization of the sensors at the work-piece surface, in ¼ diameter distance from the surface, and on the work-piece axis enabled detection of the effective strain not only in the peripheral, but also in the internal regions of the swaged-piece. Stress distribution after swaging via both the processing routes was evaluated across the perpendicular cutting plane (also depicted in Fig. 1b).

The samples were found to have similar crystal structures to that of pure tungsten [37]. In addition, a secondary phase NiCo2W was detected in this THA. In order to perform a reliable stress analysis, a high angle peak needs to be chosen to reduce the amount of error in the measurement. Therefore, the (321) plane which occurs around 131° [38] was chosen for stress calculations [39]. The crystal lengths were manually calculated from software and the step size of the XRD scan was 0.03°. The last experimental step was measurement of mechanical properties, which was performed using a Zwick machine with the strain rate of 1.3 × 10−3 s−1. 2.2. Numerical simulations

3. Results Deformation behaviour during swaging of the examined THA via both the routes, as well as the distribution of stress after swaging, was predicted using the Forge NxT commercial software. The assembly input described more in detail e.g. in [40–42] (the dies oscillate in the radial direction and rotate around the swaging axis according to selected boundary conditions) consisted of the work-piece, four swaging dies, and a clamp bar serving to feed the work-piece to the swaging head every time the dies were in the open position, i.e. after each stroke. The entire assembly was meshed using tetrahedron elements; the mesh of the work-piece having 30 mm in diameter and 100 mm in length consisted of 45 350 nodes. The initial work-piece temperature was set to 20 °C for route A, and 900 °C for route B. The dies’ initial temperature was 20 °C for both routes. Friction between the dies and work-piece was characterized with the Coulomb law (friction coefficient μ = 0.1). Material properties were set via a constitutive elasticviscoplastic model created by the software on the basis of the THA tensile test stress-strain curve. The model was determined by the Haensel–Spittel relation (Eq. (4)), where ε ̇ is equivalent strain rate, ε is equivalent strain, T is temperature, and A, m1, m2, m3, m4, m5, m7, m8, m9 are regression coefficients, the values of which were the following: A=1447.002738, m1=-0.009, m2=0.0895, m3=0.0044, m4=-0.0069, m5, m7, m8, m9 were 0.

m ̇ 3εm ̇ 8T σ = A exp(m1 T ) T m9 ε m2 exp ⎛ 4 ⎞ (1 + ε )m5 T exp(m7 ε ) ε m ⎝ ε ⎠

3.1. Surface quality Optical image of the structure of the sintered specimen surface is depicted in Fig. 3a, while Figs. 3b and c depict the surface roughness maps for the A and B swaged-pieces. Both the swaging procedures evidently affected the surface of the processed bar positively. Nevertheless, cold swaging (sample A) resulted in a smoother surface in the microscale (Fig. 3a), and also in the macroscale - Fig. 3d depicting the macroscopic photos of both the swaged rods shows the smooth swaged finish of sample A. 3.2. Physical properties The individual types of thermomechanical processing imparted certain differences in the THA’s physical properties. The homogeneousness of a material is usually demonstrated by the shear wave being half of the longitudinal wave [43]. As evident from Table 1, the longitudinal waves were less than two times the shear waves for all the samples, which suggests their heterogeneous nature. Their physical properties were found to be similar to those of pure tungsten, except for A1 and B1 samples Poisson’s ratio. The results thus suggest that the post-process heat treatment increased the heterogeneity. For example, the ratio of Vtrans to Vlong for sample A was 1.80, while for A1 it was 1.77, as can also be seen in Figs. 4a and b showing the A and A1 samples’ strain graphs. Moreover, the heat treatment evidently imparted increase in both the Young’s and Shear modulus for the samples swaged via both the routes A and B. In other words, these findings prove the essential effect of the swaging temperature on the physical properties of the processed material. The Young’s and Shear modulus were the highest for sample B1, so was the Vtrans to Vlong ratio, which indicates sample non-uniformity. However, the Poisson’s Ratio was the lowest for this sample. This was expected as sample B1 featured an

(4)

The additional processing conditions calculated with the help of the model were thermal conductivity of 173 W.m-1.K-1, specific heat of 130 J.kg-1.K-1, emissivity of 0.88, Poisson coefficient of 0.27, and density of 17.3 g.cm-3. The individual parameters were corrected based on the experimental results. The model was applicable within the temperature range of 20 – 1050 °C, strain range of 0.04 – 2, and strain rate range of 0.01 – 100 s-1. All the conditions were thus in accordance with the 3

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Fig. 3. Optical image of sintered rod surface (a); surface scan map for sample A (b); surface scan map for sample B (c); pieces swaged via routes A and B (d).

more or less homogenous. As can be seen in Fig. 4b, sample A1 exhibited the most substantial heterogeneity of all the investigated structures. This is documented also by Table 3 the values for sample A1 in which exhibited the most substantial deviations. Also, the total residual stress values within samples A1 and B1 were lower than within samples A and B. These findings confirm the essential effect of the swaging temperature on residual stress development. In order to correlate the measured residual stress values with its distribution throughout the swaged structures, detailed investigations of residual stress distribution throughout the structures were performed via internal grain misorientations analyses, the results of which are depicted in Figs. 5a to d. In the figures, the misorientations are characterized by the rainbow colour distribution in the range from 0° (blue colour) to 15° (red colour). Mutual comparison of the samples swaged via routes A and B reveals that the presence of residual stress was more significant within sample A swaged at room temperature, which also corresponds to the results of residual stress calculations summarized in Table 3. As can be seen, most of the W agglomerates featured high misorientations especially in their peripheral areas. According to XRD analyses, the distribution of stress was quite homogeneous for both the A and B samples. Figs. 5a and c also show homogeneous distribution of stress throughout the structures – the red regions depicting high misorientations were distributed more or less uniformly within the tungsten grains, and also within the γ matrix. As depicted in Figs. 5b and d, the post-process annealing at 900 °C imparted stress relaxation especially within the matrix, which introduced stress heterogeneousness throughout the structures (as shown in Figs. 4b and d) and decrease in the overall level of residual stress (as shown in Table 3).

Table 1 Calculated physical properties of THA samples based on ultrasound measurements. Sample

W

A

A1

B

B1

Young’s modulus [GPa] Shear modulus [GPa] Poisson’s ratio [-] Standard error for longitudinal wave [mm/μs] Standard error for shear wave [mm/μs] Average density [g/cm3]

340-450 130-160 0.280 -

350.52 137.18 0.278 0.0045

365.46 144.38 0.266 0.0017

359.08 141.34 0.270 0.0039

387.16 154.14 0.256 0.0071

-

0.0019

0.0034

0.0050

0.0231

17.473

17.348

17.203

17.424

oxide layer and non-uniform top surface, which resulted in the largest measured standard error in the velocity of the longitudinal and shear wave (Table 1) [33]. However, as further discussed, the XRD scan for sample B1 showed no effect of the oxide on the results. Therefore, the layer was not necessary to be removed. The results of these measurements were advantageously used to better characterize the examined material during the subsequent numerical simulation. 3.3. Residual stress The crystal lengths of the investigated samples shown in Table 2 were found to be slightly larger than of pure tungsten, however, the maximum difference was only ~ 0.001 Å. The crystal length of samples A, A1 and B increased by 0.0011 Å compared to pure tungsten, while sample B1 exhibited crystal structure very similar to pure tungsten (difference of only 0.0001 Å). The results of strain measurements performed for each sample are summarized in Figs. 4a to d; the x axis in the figures shows the sin2ψ value for each ψ angle, the regression line is depicted in blue colour. The final calculated σ11 and σ22 (see Fig. 2) residual stress values for each sample are then depicted in Table 3. The measurements show that the stress within samples A and A1 featured heterogeneities, while the stress within samples B and B1 was

3.4. Structure observations Fig. 6a depicts the original structure of the sintered and quenched samples, which were further subjected to cold and warm swaging. As can be seen, the structure featured round occasionally adjoining tungsten agglomerates surrounded with the NiCo matrix. Mutual 4

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of < 110 > parallel to the swaging direction (< 110 > || SD), whereas their peripheral regions featured orientation changes caused by substructure development. On the other hand, OIM of sample B depicted in Fig. 6c shows that the agglomerates within this sample exhibited no prevailing preferential orientation. The prevailing < 110 > || SD deformation texture within W agglomerates after the applied heat treatment demonstrates that the annealing temperature of 900 °C had no substantial effect on the structures of W agglomerates (Figs. 6c and e). The substantial presence of residual stress within sample A can also be documented by a TEM image showing the interfaces of two tungsten agglomerates and the γ matrix in between (Fig. 7a). The softer matrix exhibits discontinuities resulting from the presence of residual stress. Fig. 7b then shows a TEM scan from the peripheral location of a W agglomerate within this room-temperature-swaged sample depicting accumulated dislocations and formation of subgrains elongated in the direction of dominant imposed strain. These analyses support the above drawn conclusions on substructure development. 3.5. Numerical prediction Numerically predicted behaviour of this particular THA during a single pass of cold rotary swaging was thoroughly investigated in a previous publication [20]. The herein reported numerical study presents the results of swaging via one and two passes at room temperature and at the temperature of 900 °C. The results showed that the increased number of swaging passes, as well as the increased processing temperature, imparted evident differences; differences in deformation behaviours of the individual swaged-pieces processed via both the deformation routes were documented by the developments of effective strain in selected locations represented by the individual sensors depicted in Fig. 1 (Figs. 8c and d). When compared to single pass swaging, swaging via multiple passes imparted accumulation of the imposed strain. For route A, the maximum values of the effective strain of ~1.5 were located in the swagedpiece peripheral region, whereas the axial region featured the effective strain of approximately 0.6 after the second pass (Fig. 8a). The overall effect of the strain imposed via the two swaging passes on the entire volume of the piece swaged via route B was more substantial; the maximum effective strain values at the swaged-piece periphery reached up to the value of 2, and the axial region featured values around 1 (Fig. 8b). The results show that the differences between the effective strain values in the peripheral and axial regions of the individual swaged-pieces were comparable for both the A and B routes (1.5 and 0.6 compared to 2 and 1, respectively). Nevertheless, the effect of the imposed strain across the cross-section of the individual pieces, i.e. the effective strain development in ¼ diameter from surface, was different. Swaging via route A introduced comparable effective strain values in the axial and ¼ diameter regions after the first pass, whereas the second pass imparted certain strain heterogeneity (Fig. 8c). On the other hand, two swaging passes via route B resulted in effective strain homogenization; differences in the imposed strain distribution in the axial and ¼ diameter regions were only evident after the first pass (Fig. 8d). Considering the stress distribution across the monitored plane (Fig. 1) after the second swaging pass, swaging via route A introduced trends similar to those discussed in the previous study of a single swaging pass [30]. The peripheral and axial regions of the swaged-piece exhibited the highest values of compressive stress, whereas the interregion between these two areas featured decrease in stress values (Fig. 9a). In other words, two room temperature swaging reductions led to certain homogenization of stress across the cross-section compared to a single reduction, but the character of stress distribution was kept. The piece swaged via route B showed evident differences especially in the axial region (Fig. 9b). This treatment caused the lowest compressive stress values to occur in the axial region of the swaged-piece (contrary to route A); the highest stress values were detected in the peripheral area (likewise route A). The comparison of the predicted results and the

Fig. 4. Residual stress depicted as dependence of measured strain on sin2ψ values for samples: A (a); A1 (b); B (c); B1 (d). Table 2 Crystal structure and length for each sample. Sample

Space group

Crystal length a in Å

W pure A A1 B B1

229 229 229 229 229

3.1649 3.1660 3.1660 3.1660 3.1650

Table 3 Residual stress within investigated samples in σ11 and σ22 planes. Sample

σ11 ± error (MPa)

σ22 ± error (MPa)

A A1 B B1

-979.9 ± 20.0 -832.1 ± 71.9 -936.2 ± 14.8 -720.9 ± 18.2

-973.9 ± 20 -833.2 ± 196.3 -912.2 ± 40.3 -733.7 ± 49.7

comparison of the shapes of W agglomerates within samples A and B (Figs. 6b and d) shows that the agglomerates within sample A were deformed. As mentioned previously, the presence of misorientations within the individual W agglomerates was evident in all the samples, but the most evident in sample A. Fig. 6b, the orientation image map (OIM) for sample A in which is depicted, shows that the internal volume of the agglomerates mostly featured the single preferential orientation 5

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Fig. 5. Internal grains misorientations indicating residual stress distribution throughout structures for samples: A (a); A1 (b); B (c); B1 (d).

exhibited rapid fracture without significant elongation.

XRD results of residual stress measurements show a good correlation; experimental sample A exhibited the highest values of compressive stress of almost 1 000 MPa, whereas the maximum compressive stress values within sample B were only slightly higher than 900 MPa (Section 3.2.).

4. Discussion Swaging generally affects the entire volume of the processed material. However, given by the character of the method, the absolute value of the imposed strain is the highest on the swaged-piece periphery and decreases gradually across the cross-section towards the axis. Correspondingly to the two components of the affecting swaging force [21], the plastic flow during RS consists of the axial and tangential component. The vector profiles for both the swaging routes (not shown here but studied previously in [20]) revealed that the axial plastic flow component was dominant for both the herein introduced A and B routes. Nevertheless, the increase in swaging temperature imparted intensification of the tangential component for swaged-piece B, the plastic flow of which was facilitated significantly. This fact was documented not only by the evidently higher values of the predicted effective strain, but also by the elongated shapes of the tungsten agglomerates (as seen in Figs. 5 and 6) and the characteristic shape of the swaged-piece ends; the intensive material flow in the axial region of the swaged-piece processed via route B resulted in visible bulging (Figs. 3d, 8b). However, RS influences not only the shape of the final swagedpiece, but also the individual structural features. Mutual comparison of the shapes of W agglomerates within samples A and B (Figs. 6b and 5d) shows that the agglomerates within sample A had deformed shapes. On the other hand, the shapes of the agglomerates within sample B were almost equiaxed (close to the round shape), i.e. their shapes were similar to the shapes of the agglomerates right after sintering (Fig. 6a). Given by the very high stacking fault energy of nickel, the matrix is easy to be deformed at the temperature of 900 °C [44]. Katavič et al. [45] previously reported the hardening rate of γ matrix during intensive

3.6. Mechanical properties The stress-strain curves for the sintered material, swaged-pieces processed via routes A and B by one and two passes, and heat treated A1 and B1 samples are depicted in Fig. 10, while the values of characteristic mechanical properties for the samples are summarized in Table 4. As can be seen, the lowest ultimate tensile strength (UTS) of approximately 860 MPa was recorded for the sintered (initial) material. On the other hand, this material state exhibited relatively high elongation to failure, more than 18 %, which was the highest value of all the samples, except route B first pass (see Table 4). The strength increased significantly for all the swaged-pieces; increasing imposed strain generally introduced increased strength and decreased plasticity (elongation) for both the swaging routes. The overall strengthening was the highest for samples A, whereas the increased temperature resulted in relatively high plasticity for samples B. The highest elongation to failure of almost 24% was recorded for the piece swaged at 900 °C via a single pass. Mutual comparison of the curves for routes A and B shows that the increased swaging temperature imparted a more gradual strengthening especially during the first swaging pass, i.e. the slope of the curves for pieces swaged at 900 °C (route B) was smoother than for pieces swaged at room temperature (route A). The subsequent heat treatment (samples A1 and B1) introduced decrease in plastic properties, however, the strength did not change drastically. As can be seen in Fig. 10, the structure changes introduced by the heat treatment primarily resulted in a greater steepness of the final curves, i.e. the heat-treated samples 6

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Fig. 6. SEM-BSE image of sintered and quenched structure (a); orientation image maps showing grains orientations for samples: A (b); A1 (c); B (d); B1 (e).

Fig. 7. Transmission electron microscope images for sample A: W agglomerates/matrix interfaces (a); substructure within W agglomerate (b).

7

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Fig. 8. Distribution of effective strain across a longitudinal cutting plane during 2nd swaging pass: route A (a); route B (b). Effective strain in monitored locations: route A (c); route B (d).

plastic deformation of THAs to be almost two times higher than the hardening rate of W agglomerates at the deformation of approximately 5 %. The rate then increases to the deformation of approximately 15 % and then saturates. The herein observed differences in the structures and mechanical behaviours of swaged-pieces A and B can be explained as follows: swaging via route A resulted in strengthening and exhaustion of plasticity within the γ matrix, which was also documented by Fig. 6b showing tungsten agglomerates deformed by the effect of the imposed strain, which was sufficient for the matrix to harden enough to further transfer the strain to the agglomerates, and Fig. 7a, the exhaustion of plasticity within the matrix in which was evident. On the other hand, swaging at the elevated temperature enabled the matrix to dynamically recrystallize and provided it with the ability to consume a greater amount of the imposed strain. This was documented by Fig. 6d the round-shaped agglomerates in which were depicted. The effects of the individual swaging routes on the interactions between the harder agglomerates and softer matrix also affected nonnegligibly the stress-strain behaviour of the swaged-pieces. The higher swaging temperature facilitating the plastic flow of the swaged-piece

Fig. 10. Experimental stress-strain curves for sintered, swaged, and heat treated material states.

contributed to higher homogeneity of strain distribution and consequent decrease in residual stress for sample B. On the other hand, the plastic flow of the material in the vicinity of the swaging dies (i.e. the

Fig. 9. Numerically predicted stress distribution after second pass: route A (a); route B (b). 8

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5. Conclusions

Table 4 Mechanical properties of THAs after swaging and heat treatment. State

Initial state Route A Route A Route A + HT Route B Route B Route B + HT

Deformation degree

Yield strength

Tensile strength

Yield ratio

Elongation

(-)

(MPa)

(MPa)

(-)

(%)

-

690

859

0.803

18.3

0.50 0.81 0.81

1789 1738

1409 1868 1754

0.957 0.991

9.2 5.6 3.8

0.50 0.81 0.81

930 1503

1013 1643 1571

0.918 0.956

23.8 7.8 2.9

This work numerically and experimentally investigated the effects of two different thermomechanical treatments on pre-sintered WNiCo rods. Processing was performed via rotary swaging at 20 °C and at 900 °C, the temperatures were selected with the aim to minimise possible grain growth and oxidation to which tungsten in highly prone. Both the processing regimes resulted in increase in strength, the highest UTS of 1868 MPa was acquired after two passes of cold swaging, while the highest plasticity of 23.8 % was recorded after a single swaging pass at 900 °C. The increased plasticity can primarily be attributed to structure relaxation and decrease in residual stress within the matrix. Both the experimental and numerically predicted results revealed that swaging at room temperature introduced the highest compressive residual stress occurring especially within the γ matrix and W grains peripheries, however, the values decreased substantially after post-process heat treatment at 900 °C (but annealing introduced stress heterogeneity). On the other hand, two swaging passes at 900 °C resulted in a more or less homogeneous distribution of the effective strain across the swaged-piece cross-section (except the peripheral region influenced by the swaging process the most intensively), and homogenization of the distribution of residual stress. Suffice to say, both the proposed swaging routes and their combinations with post-process heat treatments resulted in successful fabrication of high-strength THA rods suitable for subsequent product manufacturing; no substantial microstructure instabilities were observed.

peripheral and sub-peripheral regions) was aggravated by the cooling effect of the dies and direct influence of the swaging force tangential component. By these reasons, these regions featured substantially higher values of compressive stress and effective strain. The residual stress values acquired via XRD analyses showed very good correlation with the numerically predicted results. Swaging via route A imparted slightly higher integral values of residual stress than swaging via route B, which is evident from table 3 and also from Figs. 9a and b. The possibilities of comparison of the herein presented results with studies by others are limited. Nevertheless, considering the conclusions drawn in the previous study [30], the increasing number of room temperature swaging reductions imparts certain homogenization of residual stress across the cross-section of the swaged-piece. The most substantial strain heterogeneity was detected for sample A1. This can primarily be attributed to structure changes which were imparted by the intensive plastic deformation and preserved due to the aggravated restoration; the post-process heat treatment did not introduce sufficient energy to fully relax the deformed structure. Also, the applied heat treatment introduced mutual diffusion of W and Ni/Co phases, which resulted in dramatic decrease in the plastic properties of the heat treated samples; similar diffusion effect was observed also for sample B1. The total residual stress within samples A1 and B1 was lower than within samples A and B, which points to the positive influence of the increased temperature on structure relaxation, especially within the γ matrix. After cold swaging, the matrix exhibited typical FCC deformation texture - < 111 > || SD. The swaging/annealing temperature of 900 °C introduced recrystallization and consequent grain refinement within the matrix, which reduced the intensity of the preferential grains orientation. As regards W agglomerates, the typical < 110 > fibre orientation occurring within deformed BCC metals exhibits the tendency to change to < 111 > orientation after recrystallization [44]. The prevailing < 110 > || SD orientation within the agglomerates demonstrates that nor the swaging/annealing temperature of 900 °C had any substantial effect on the structures of the agglomerates; softening was thus provided especially by the γ matrix, the recrystallization temperature for which is significantly lower than for the tungsten agglomerates [46]. Considering the fact that the strain imposed to sample B was consumed primarily by the matrix, the observed final orientations of W agglomerates were most probably affected not only by the swaging process, but also by the previous sintering, the texture developed during which could have been preserved i.e. the properties and structure of the material after sintering, can non-negligibly influence the final performance of the THA product. The influence of sintering temperature on the deformation behaviour and fracture mechanism of THA was numerically and experimentally investigated by others (e.g. [47]). For the herein presented material, more detailed texture and grain size analyses are planned to be performed in our following study.

Data availability statement The raw data required to reproduce this study cannot be shared at this time as the data are part of an ongoing study. Acknowledgments The authors acknowledge the support of 19-15479S Project of the Grant Agency of the Czech Republic and the MACH1 COMET project at Swansea University funded by the Welsh Government. We would like to thank Advance Imagining of Materials (AIM) for giving access to the XRD and Tom Dunlop for his support with XRD, as well as Dr. Petr Král for his help with TEM and EBSD data evaluation. We also wish to acknowledge the use of the EPSRC hosted by the Royal Society of Chemistry References [1] Y. Li, K. Hu, X. Li, X. Ai, S. Qu, Fine-grained 93W–5.6Ni–1.4Fe heavy alloys with enhanced performance prepared by spark plasma sintering, Mater. Sci. Eng. A 573 (2013) 245–252, https://doi.org/10.1016/j.msea.2013.02.069. [2] B. Katavic, M. Nikacevic, Z. Odanovic, Effect of cold swaging and heat treatment on properties of the P/M 91W-6Ni-3Co heavy alloy, Sci. Sinter. 40 (2008) 319–331, https://doi.org/10.2298/SOS0803319K. [3] U. Ravi Kiran, A. Sambasiva Rao, M. Sankaranarayana, T.K. Nandy, Swaging and heat treatment studies on sintered 90W-6Ni-2Fe-2Co tungsten heavy alloy, Int. J. Refract. Met. Hard Mater. 33 (2012) 113–121, https://doi.org/10.1016/j.ijrmhm. 2012.03.003. [4] A. Ghaderi Hamidi, H. Arabi, J. Vahdati Khaki, Sintering of a nano-crystalline tungsten heavy alloy powder, Int. J. Refract. Met. Hard Mater. 80 (2019) 204–209, https://doi.org/10.1016/J.IJRMHM.2019.01.016. [5] L. Xu, F. Xiao, S. Wei, Y. Zhou, K. Pan, X. Li, J. Li, W. Liu, Development of tungsten heavy alloy reinforced by cubic zirconia through liquid-liquid doping and mechanical alloying methods, Int. J. Refract. Met. Hard Mater. 78 (2019) 1–8, https:// doi.org/10.1016/J.IJRMHM.2018.08.009. [6] S. Marschnigg, C. Gierl-Mayer, H. Danninger, T. Weirather, T. Granzer, P. Zobl, Non-destructive measurement of the tungsten content in the binder phase of tungsten heavy alloys, Int. J. Refract. Met. Hard Mater. 73 (2018) 215–220, https:// doi.org/10.1016/J.IJRMHM.2018.02.005. [7] L. Kunčická, T.C. Lowe, C.F. Davis, R. Kocich, M. Pohludka, Synthesis of an Al/ Al2O3 composite by severe plastic deformation, Mater. Sci. Eng. A 646 (2015) 234–241, https://doi.org/10.1016/j.msea.2015.08.075. [8] Q. Wei, L.J. Kecskes, Effect of low-temperature rolling on the tensile behavior of

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