In order to examine the effects of the sintering temperature on chemical composition, i.e., distribution of the individual elements within the structures, SEM-EDX analyses were performed for selected sintered pieces. Figure 3a shows the chemical composition for the sample sintered at 1500 ◦C for 30 min depicted via distributions of the individual elements within the scanned area. Such distribution was typical for the samples sintered at the temperature of 1500 ◦C and lower. Figure 3c then shows quantitative evaluation of the individual elements within the scanned map. The chemical composition depicted via distributions of the individual elements within the sample sintered at 1525 ◦C for 30 min is shown in Figure 3b, corresponding quantitative evaluation of the individual elements within the scanned area of the sample is then depicted in Figure 3d. Such distribution of the individual

elements was typical for the samples sintered at and above the temperature of 1525 ◦C. The maps in Figure 3b show that tungsten is primarily concentrated in the agglomerates surrounded by the NiCo matrix.

**Figure 2.** Structure scans of individual sintered pieces: 1450 ◦C/30 min (**a**); 1475 ◦C/30 min (**b**); 1500 ◦C/30 min (**c**); 1525 ◦C/30 min (**d**); 1550 ◦C/30 min (**e**). In OM images, white/light areas depict the NiCo matrix, grey areas depict tungsten, and black areas depict voids/cracks.

**Figure 3.** SEM-EDX maps of distribution of individual elements (W-red, Ni-blue, Co-violet) for sintered pieces: 1500 ◦C/30 min (**a**); 1525 ◦C/30 min (**b**); quantitative evaluation of individual elements within maps depicted in (**a**) and (**b**) in investigated structures, in wt. %: 1500 ◦C/30 min (**c**); 1525 ◦C/30 min (**d**).

#### *3.2. E*ff*ect of Sintering Time on Structures*

The subsequent research step was the investigation of the effects of variable sintering time at the critical sintering temperature of 1500 ◦C. The results of the analyses are depicted in Figure 4a–e, the SEM scans of samples sintered for 60, 90, 120, 150, and 180 min in which are depicted, respectively. Evidently, increasing the sintering time at the temperature of 1500 ◦C to at least 60 min leads to the formation of the microstructures consisting of tungsten agglomerates surrounded by the NiCo matrix.

Despite the fact that increasing the sintering time has favorable effects on the structure character when examined via SEM, more detailed TEM investigations showed that increasing the time dwell on the high temperature leads to inter-diffusion of tungsten into the NiCo matrix, as documented by Figure 5a,b. Figure 5a shows a TEM scan of a W-NiCo-W interface within the sample sintered at 1500 ◦C for 180 min, Figure 5b then shows the chemical composition (in wt. %) along a line scan measured across the mentioned interface. As can be seen, the content of tungsten increased to about 50 wt. % in the inter-agglomerate region by the effect of increased sintering time (negligible tungsten content within the NiCo matrix was observed for the sample sintered at 1500 ◦C for 60 min (not shown here)).

**Figure 4.** Structure scans of individual pieces sintered at 1500 ◦C for: 60 min (**a**); 90 min (**b**); 120 min (**c**); 150 min (**d**); 180 min (**e**).

**Figure 5.** TEM scan showing a W-NiCo-W interface and measured line scan area for sample sintered at 1500 ◦C for 180 min (**a**); chemical composition of measured line scan (**b**).

#### *3.3. E*ff*ect of Quenching on Structures*

Investigation of the effects of quenching into room temperature water right after sintering was performed for all the samples sintered in the temperature range from 1450 to 1550 ◦C for 30 min, as well as the sample sintered at 1500 ◦C for 180 min. As shown by the structure analyses the results of which are depicted in Figure 6a–f, quenching does not have any significant effect on the structure character, as the quenched structures were similar to the non-quenched ones; quenched samples sintered at/above the temperature of 1525 ◦C for 30 min, and the sample sintered at 1500 ◦C for 180 min, exhibited the structure consisting of W agglomerates and the NiCo matrix, while the other analyzed samples featured more or less randomly distributed W/NiCo locations.

**Figure 6.** Structure scans of individual quenched pieces sintered at: 1450 ◦C/30 min (**a**); 1475 ◦C/30 min (**b**); 1500 ◦C/30 min (**c**); 1500 ◦C/180 min (**d**); 1525 ◦C/30 min (**e**); 1550 ◦C/30 min (**f**).

#### *3.4. Mechanical Properties*

Microhardness was measured for all the sintered pieces in order to analyze the effects of the sintering temperatures on their properties. The results of the measurements are summarized in Table 1. As regards to the samples sintered for 30 min, the highest average microhardness value was recorded

for the sample sintered at 1550 ◦C. However, the highest microhardness value of all was measured for the sample sintered at 1500 ◦C for 180 min. This finding led to the performance of a supplementary experiment, in which a sample was sintered at the temperature of 1525 ◦C for 180 min. However, for the elevated temperature, the effect of increasing sintering time is not positive, as the microhardness decreased dramatically when compared to the sample sintered at 1500 ◦C for 180 min.

**Table 1.** Average microhardness for sintered pieces.

Last but not least, the tensile tests were performed for the sintered and subsequently quenched pieces; the resulting stress–strain curves are depicted in Figure 7. Ass can be seen from the results, the samples sintered at/below 1500 ◦C for 30 min, featuring structures with randomly distributed W/NiCo locations and voids, exhibit very low plasticity (lower than 2.5%). The 1450 and 1757 ◦C samples exhibited the lowest plasticity of less than 1%. On the other hand, structures featuring the homogeneous distribution of tungsten agglomerates within a NiCo matrix exhibit a very high strength of more than 1000 MPa, together with excellent plasticity (elongation to failure exceeding 20%).

**Figure 7.** Stress–strain curves for sintered and subsequently quenched pieces.

#### **4. Discussion**

Microstructure observations showed that, during sintering at temperatures above the critical value of 1525 ◦C, the original powder particles consolidated and the structures exhibited the desired distribution of tungsten agglomerates with the sizes of several dozens of micrometres surrounded by a matrix primarily consisting of Ni and Co, whereas the samples sintered below this temperature featured more or less randomly distributed locations with high tungsten or nickel-cobalt contents. Moreover, the structures of pieces sintered below 1525 ◦C also featured voids and gaps. Sintering at low temperatures thus does not impart the desired structure homogenization and spheroidization of tungsten powder particles into agglomerates. In other words, sintering at temperatures at/above 1525 ◦C leads to local melting of the alloying elements, the melting temperatures of which are lower than the sintering one (1455 ◦C for nickel and 1495 ◦C for cobalt [28]), which supports the development of the desired structure character via diffusion. By this reason, especially the samples sintered below 1500 ◦C exhibited the presence of voids/gaps, as documented in Figure 2a,b.

Having performed the first part of the study, we found that the optimum sintering regime among the investigated ones, from the viewpoint of achievement of the required structure character, is the temperature of 1525 ◦C and time of 30 min. However, as the sintering temperature is definitely higher than the melting temperatures of both the alloying elements, the work pieces exhibit the tendency to gravity sedimentation during sintering, which increases the cross-sectional ovality. In other words, increasing the sintering temperature causes the shape of the cross-section of the sintered pieces to be oval, more than circular. By this reason, an additional experiment with increasing the sintering time at the sintering temperature of 1500 ◦C was performed.

Generally, increasing the sintering time was found to affect the mean size of the tungsten agglomerates—it increased with increasing sintering time [29], as well as to support inter-diffusion of the individual elements. Previous neutron powder diffraction measurements identified two individual phases in the sintered pieces featuring the required structure character, the main one of which was α-W featuring the B2 structure (W-B2 phase) [29]. The second phase, NiCo2W, with the weight fraction of 6–7% had a pure Ni-like structure (FCC) with the lattice parameter of about 3.60 Å, which is, however, slightly larger than for pure nickel (3.55 Å). This fact indicated alloying of the matrix with larger atoms (atomic radius of W, Ni, and Co is 139, 124, and 125 pm, respectively), and thus revealed the presence of a certain amount of diffused W in the NiCo-based matrix. This phenomenon was explained by Chuvil'deev et al. [30], who documented that introducing distortions to the tungsten lattice, which can be performed by introducing energy via high sintering temperature and/or time, or via subsequent plastic deformation, leads to formation of strong fields of internal stress, which decreases the activation energy necessary for decoupling of tungsten atoms and supports their diffusion to the NiCo matrix. These phenomena are behind the increase in microhardness for the 1500 ◦C/180 min sintered piece. On the other hand, the 1525 ◦C/180 min sintered piece exhibited rapid decrease in microhardness when compared to the 1500 ◦C/180 min one. This can be explained via softening processes occurring within the matrix [17].

The subsequent quenching does not affect the structure character of the sintered pieces. However, it has positive effects on the mechanical properties, testing of which revealed excellent plasticity exceeding 20%, together with the ultimate tensile strength (UTS) of more than 1000 MPa for the pieces sintered at 1525 and 1550 ◦C for 30 min and quenched. The fact that the UTS of the 1550 ◦C sintered piece is slightly lower than of the 1525 ◦C one can be attributed to the occurring restoration phenomena supported by the increased sintering temperature (further temperature increase from 1525 ◦C introduces similar effects as increasing the sintering time at 1525 ◦C) [17,28].

#### **5. Conclusions**

The study focused on the investigations of the effects of sintering temperature, sintering time, and optional subsequent quenching in water on the structures and mechanical properties of a 93W6Ni1Co tungsten heavy alloy. The results revealed the following:

The optimized preparation regime providing the 93W6Ni1Co alloy with structure and properties being the most favorable for subsequent processing is sintering at 1525 ◦C for 30 min and subsequent quenching in water.

**Author Contributions:** The contributions of the individual authors were the following: methodology, R.K., L.K. and Z.K.; experimental validation, R.K., L.K.; microscopy investigation and evaluation, L.K.; writing-original draft preparation, L.K., R.K.; writing—review and editing, L.K.; project administration, Z.K.; funding acquisition, R.K., Z.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** The authors acknowledge the financial support of the FV40286 TRIO programme by ministry of Industry and Trade, and the 19-15479S Project of the Grant Agency of the Czech Republic.

**Acknowledgments:** The experiments were performed using equipment available at the Technical University of Ostrava. We also acknowledge the help of IPMinfra research infrastructure and working team of Institute of Physics of Materials, CAS.

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

#### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **A**ff**ecting Structure Characteristics of Rotary Swaged Tungsten Heavy Alloy Via Variable Deformation Temperature**

#### **Adéla Macháckov ˇ á, Ludmila Krátká \*, Rudolf Petrmichl, Lenka Kunˇcická and Radim Kocich \***

Faculty of Materials Science and Technology, VŠB–Technical University of Ostrava, 17. listopadu 15, 708 00 Ostrava, Czech Republic<EMAIL_ADDRESS>(A.M<EMAIL_ADDRESS>(R.P<EMAIL_ADDRESS>(L.K.)

**\*** Correspondence<EMAIL_ADDRESS>(L.K<EMAIL_ADDRESS>(R.K.); Tel.: +420-596-99-44-33 (L.K.)

Received: 2 November 2019; Accepted: 12 December 2019; Published: 13 December 2019

**Abstract:** This study focuses on numerical prediction and experimental investigation of deformation behaviour of a tungsten heavy alloy prepared via powder metallurgy and subsequent cold (20 ◦C) and warm (900 ◦C) rotary swaging. Special emphasis was placed on the prediction of the effects of the applied induction heating. As shown by the results, the predicted material behaviour was in good correlation with the real experiment. The differences in the plastic flow during cold and warm swaging imparted differences in structural development and the occurrence of residual stress. Both the swaged pieces exhibited the presence of residual stress in the peripheries of W agglomerates. However, the NiCO matrix of the warm-swaged piece also exhibited the presence of residual stress, and it also featured regions with increased W content. Testing of mechanical properties revealed the ultimate tensile strength of the swaged pieces to be approximately twice as high as of the sintered piece (860 MPa compared to 1650 MPa and 1828 MPa after warm and cold swaging, respectively).

**Keywords:** tungsten heavy alloy; rotary swaging; finite element analysis; deformation behaviour; residual stress

#### **1. Introduction**

For their exceptional mechanical (high strength) and physical (high density, melting point, etc.) properties, tungsten heavy alloys (THAs) are primarily used to shield radiation or to block kinetic energy [1,2]. Nevertheless, they can also be used advantageously for other demanding applications, such as for production of therapeutic devices in oncology, for kinetic penetrators in the military, or as aircraft counterbalances [3,4]. THAs are typically produced via powder metallurgy—they are isostatically pressed and subsequently sintered at temperatures between 1000 and 1500 ◦C—from powder mixtures containing a majority (over 90 wt.%) of tungsten, in addition to other alloying elements featuring lower melting temperatures (e.g., Ni, Co, Fe), the boundaries of the grains of which melt during sintering and ensure binding of the tungsten particles and elimination of porosity via diffusion [5].

To enhance the strength of THAs, the tungsten content can be increased to 98 wt.%; however, the increase in strength is typically at the expense of plasticity. Additions of alloying (i.e., binding) elements generally decrease the strength, but increase plasticity. Contents of the alloying elements (i.e., matrix) lower than 3 wt.% can cause brittleness of the final product [6]. On the other hand, high contents of alloying elements contribute to inhomogeneity of the mechanical properties and usually also to an uneven shape of the sintered piece cross-section due to gravity sedimentation during sintering. Under optimized conditions, the matrix is homogenously distributed in the gaps between the tungsten agglomerates [7].

In addition to modification/optimization of their chemical composition, the mechanical properties of THAs can also be enhanced by post-sintering deformation processing, which can advantageously be performed via methods of severe plastic deformation [8–10]. Technologies applying intensive plastic deformation have recently gained attention primarily due to their ability to effectively refine the structural units and consequently enhance the mechanical and utility properties of the processed materials. Severe plastic deformation (SPD) methods, such as equal channel angular pressing (ECAP) [11–13], twist channel (multi) angular pressing (TCAP, TCMAP) [14,15], or high pressure torsion (HPT) [16], have been proven to be very effective; however, they are discontinuous and are only applicable for limited volumes of material. On the other hand, methods such as ECAP-conform [17,18], accumulative roll bonding (ARB) [19,20], or rotary swaging (RS) [21,22] are continuous and can favourably be used to manufacture relatively large (semi)products. Given its versatility, the RS technology is applicable in various industrial branches.

Rotary swaging is an incremental process, i.e., it is characterized by gradual increments of the imposed strain provided by high-frequency strokes of the swaging dies. The method can be applied under cold, as well as warm and hot conditions and is used to process solid work pieces, but also to produce hollow tubes and shaped axisymmetric products [23–25]. The processed materials are typically conventionally cast alloys, such as magnesium [26] and aluminium alloys [27]. Nevertheless, the favourable stress state and incremental character enables the method to advantageously be applied also to process composites, the example of which can be mechanically bonded Cu/Al clad composites for electrotechnic applications [28,29], powder-based materials, demanding compounds based on zirconium [30], or tungsten heavy alloys [31].

The aim of this work was to investigate material behaviour of the studied WNiCo tungsten heavy alloy during cold and warm rotary swaging, and the effects of the individual processing steps on its structure and properties. To better characterize the influence of the used technology, especially the effect of the applied pre-swaging induction heating, experimental investigations were supplemented with finite element analyses. The predicted results were verified by real swaging experiments. Attention was also given to characterization of residual stress within the structure.

#### **2. Materials and Methods**

#### *2.1. Numerical Prediction*

The first step of the performed study was the numerical prediction of the applied induction heating process, which was performed via Finite Element Method (FEM) using FORGE (NxT version, Transvalor, Sophia Antipolis, France) software. The used computational model was a model based on the Maxwell's equations available in the software. The prediction was implemented primarily to optimize control of the heating process and determine the heating time necessary to achieve homogeneous temperature distribution throughout the work piece cross-section before the warm rotary swaging, i.e., to predict behaviour of the studied THA during induction heating. The investigated work piece, as well as the used inductor, were modelled according to the experimentally used pieces (Figure 1a).

The simulation was designed considering two main aims. Firstly, minimisation of the computational time. For this reason, meshing was performed thoughtfully to ensure both reliable results and reasonable computational time. The mesh of the work piece consisted of 2700 volume tetrahedral elements, whereas the mesh of the inductor consisted of 21,314 nodes in total (Figure 1b).

The second aim was minimisation of the real heating time. For this reason, the simulation (and the following experiment) involved two-step heating. The first step, defined with the current of 68 A and frequency of 16.5 kHz, was designed to provide the quickest possible achievement of the swaging temperature of 900 ◦C, whereas the second step ensured homogenisation of the temperature throughout the cross-section, i.e., minimised temperature gradient from the surface to the axis of the work piece. Once the surface temperature of 900 ◦C had been reached, the second heating step, defined by a current of 50 A and a frequency of 15.8 kHz, began. However, possible development of residual stress resulting from the variable heating rate and occurring temperature gradients was also considered. The entire heating process was optimized for the maximum temperature of 900 ◦C not to be exceeded by more than 20 ◦C. To ensure thorough evaluation and control of the process, three locations at the axial longitudinal cut through the work piece, depicted as sensors 1, 2, and 3 (Figure 1c), were selected and monitored during the entire heating process.

**Figure 1.** Schematic depiction of setup of induction heating, meshed (**a**); inductor with mesh (**b**); axial longitudinal cut through work piece with depicted monitored sensors 1, 2, and 3 (**c**).

The FORGE NxT software was also used to perform numerical prediction of deformation behaviour of the investigated 93W6Ni1Co alloy. The conditions for the simulation were selected to correspond to the real ones, and the entire simulation consisted of two simulation steps, corresponding to the real two swaging passes on which were modelled. The assembly for the simulation consisted of four swaging dies and the work piece, which was fed continuously towards the swaging head with the dies (Figure 2). The dies were created as shells, while the work piece was designed as a full deformable body with the mesh element size of 0.3 mm. Deformation behaviour of the material was characterized via the Haensel-Spittel equation, presented as Equation (1),

$$\sigma = A \exp(m\_1 T) T^{m\_9} \varepsilon^{m\_2} \exp\left(\frac{m\_4}{\varepsilon}\right) (1 + \varepsilon)^{m\_5 T} \exp(m \gamma \varepsilon) \varepsilon^{m\_3} \varepsilon^{m\_8 T},\tag{1}$$

where . ε is the equivalent strain rate (s−1), ε is the equivalent strain (-), T is the temperature (◦C), and A, m1, m2, m3, m4, m5, m7, m8, and m9 are regression coefficients, the values of which were set based on regression analyses of plastometric test results (A1 = 1447.002738, m1 = −0.009, m2 = 0.0895, m3 = 0.0044, m4 = −0.0069, m5, m7, m8, and m9 were equal to 0). The dimensions of the test specimens for the plastometric tests were 10 mm in diameter and 150 mm in length. Friction was defined via the Coulomb law, μ = 0,1, and the specific parameters for the simulation were the following: Poisson coefficient 0.3, specific heat 130 J·kg<sup>−</sup>1K−1, density 18.5 g·cm<sup>−</sup>3, and thermal conductivity (λ<sup>20</sup> <sup>=</sup> 149.3 W·m<sup>−</sup>1K−1, <sup>λ</sup><sup>900</sup> <sup>=</sup> 85.7 W·m<sup>−</sup>1K−1).

**Figure 2.** Schematic depiction of setup of rotary swaging.

#### *2.2. Experimental Verification*

The selected 93W6Ni1Co pseudo-alloy with the chemical composition of 92.6 wt.% W, 5 wt.% Ni and 2.40 wt.% Co (determined by SEM-EDX analysis) was prepared from powders with the mean grain size of 2.78 μm, containing ~<13 ppm of Fe, Mo, Cr, Al, Ca and other impurities. The process involved mixing of powders, cold isostatic pressing at 400 MPa, and subsequent sintering at 1525 ◦C for 20 min under a protective atmosphere followed by quenching in water. The protective atmosphere is hydrogen during sintering and argon during cooling. This preparation procedure was performed at ÚJP Praha a.s.

The sintered pieces were subsequently processed via rotary swaging, the initial temperatures for which were 20 ◦C (cold swaging) and 900 ◦C (warm swaging). The swaging temperatures were selected based on our previous experimental studies [3,31]. The room temperature was selected in order to maximize the possible achievable strength and to investigate, whether the material would maintain reasonable plasticity. The temperature of 900 ◦C was selected as the highest applicable temperature, since our previous studies showed that processing at temperatures higher than 900 ◦C resulted in massive oxidation causing rapid embrittlement during processing of the WNiCo. The temperature control during induction heating was provided by optical pyrometer (surface temperature), and a couple of thermocouples of type K (one on the surface, another in the work piece axial region). The initial diameter of the sintered pieces was 30 mm, the length was 100 mm. Both were gradually swaged down in two consequent reduction steps to the final diameter of 20 mm. Such swaged pieces are suitable, e.g., for the fabrication of kinetic penetrators [26]. The following structural analyses performed on cross-sectional cuts of the swaged pieces focused primarily on the influence of the applied deformation temperature on (sub)structure development and residual stress. The observations were performed via scanning electron microscopy (SEM EBSD analyses). The cut samples were mechanically ground on SiC papers and subsequently polished using Eposil F substance (Saphir 520 device, ATM, Germany). The EBSD analyses were performed in the sub-surface sample region, 1 mm from the outer rim of the swaged-rod, with the scan step of 0.25 μm using a Tescan Lyra 3 equipment (TESCAN Brno s.r.o, Brno, CZ) with a NordlysNano EBSD detector (Oxford Instruments, Abingdon-on-Thames, UK). The scans were evaluated using the ATEX (Win10 version) software [32]. Residual stress was characterized via analyses of internal grains misorientations in the scale from 0◦ to 15◦ (rainbow colour distribution). TEM images were acquired on ion polished thin foils with a JEOL 2100F (JEOL, Akishima, Tokio prefecture, Japan) device. The tensile tests were performed with the strain rate of 10−<sup>3</sup> using a Testometric M500-50CT testing machine (Testometric Co. Ltd., UK) on 150 mm long testing samples with circular cross-sections.

#### **3. Results**

#### *3.1. FE Analyses*

#### 3.1.1. Induction Heating

As shown by the FEA results, induction heating imparted the development of zones featuring various temperatures, similar to conventional heating in furnaces. As documented by Figure 3a, depicting the temperature field throughout the work piece in the moment of achieving the surface temperature of 900 ◦C (heating time 12 s), the difference between the temperatures in the surface and axial work piece regions was up to 350 ◦C in this moment. Nevertheless, the temperature gradient decreased rapidly as heating continued. The results showed that the heating time necessary to heat the THA work piece to the required temperature of 900 ◦C with the maximum allowable deviation of 20 ◦C and, at the same time, to provide homogeneous temperature distribution throughout the entire work piece, was 24 s (Figure 3b,c).

**Figure 3.** Temperature field during induction heating in time: 12 s (**a**); 24 s (**b**); distribution of temperature in monitored sensors (**c**).

The temperature distribution depicted in Figure 3b shows that the temperature field across the cross-section was not homogeneous along the work piece, especially during the second heating step. This is also confirmed by Figure 3c, which shows the differences between the temperatures in the individually monitored sensors. This phenomenon was a result of the geometry of the used inductor. The temperature increase was the highest in the surface work piece region up to the heating time of 10 s; the heating rate was lower in the other two monitored locations during that time. The lowest heating rate was monitored in the axial work piece area. The second heating step then imparted rather unexpected material behaviour, the heating rate in the peripheral work piece area decreased, while the other two monitored areas kept more or less the same heating rate. This behaviour finally introduced a decrease in the temperature gradient, i.e., homogenisation of the temperature field throughout the work piece cross-section, which is documented by Figure 3c.

#### 3.1.2. Deformation Behaviour

The predicted developments of temperature in the surface and axial regions during processing of both the work pieces, swaged at 20 ◦C and 900 ◦C, are depicted in Figure 4a,b, respectively. As shown in the Figures, the initial temperatures for both the swaged pieces were higher than the original swaging temperatures due to the previous deformation history (first swaging pass).

**Figure 4.** Temperature developments in surface and axial regions of the cold-swaged piece (**a**) and the warm-swaged piece (**b**).

As can be seen from Figure 4a, the temperature in the surface area increased gradually to almost 200 ◦C during cold swaging, whereas the temperature in the axial area remained more or less constant. On the other hand, the warm-swaged piece exhibited a decrease in temperature during swaging; the temperature decrease rate was comparable for both the surface and the axial areas (Figure 4b). However, multiple affecting phenomena need to be considered for the surface areas of the swaged pieces. The surface temperature increases due to the effect of friction with the rotating dies being in contact with the surface. On the other hand, the dies are cooler than the swaged piece, which causes a certain portion of the generated heat to dissipate via the swaging dies.

Figure 5a depicts the predicted developments of the imposed effective strain in the surface and axial regions of both the swaged pieces. The comparison of both the curves for the cold-swaged piece shows that the imposed strain was substantially higher in the surface region of the swaged piece, where it reached a maximum value of almost 4. The figure also shows that the maximum effective strain in the axial region of the cold-swaged piece reached to the value of 1.2. On the other hand, warm swaging imparted an increase in the imposed effective strain, which reached and even slightly exceeded the value of 2, approximately in the middle of the swaging pass. Mutual comparison of the effective strain developments for both the swaged pieces reveals that the strain gradient from the surface towards the swaged piece axis was substantially lower for the warm-swaged piece. Nevertheless, the maximum imposed effective strain was comparable for both.

**Figure 5.** Developments of imposed effective strain in surface and axial regions of both swaged pieces (**a**); effective strain across cross-section of warm-swaged piece (**b**).

The distribution of the effective strain across the cross-sectional cut through the warm-swaged piece is depicted in Figure 5b. The figure shows that the imposed effective strain reached a value of 4 at the very periphery of the swaged piece. The imposed strain then gradually decreased towards the axial region where it reached the values of approximately 2.

Figure 6a,b depicts the material flow during swaging of the cold and warm swaged piece, respectively. Given by the incremental nature of the swaging process and rotary movement of the swaging dies, material flow during swaging is variable and quite complex. The material flow was different for both the swaged pieces, and also changed its character during the individual swaging pass. At the beginning of swaging, the axial material flow dominated for both the cold and warm swaged pieces; however, the plastic flow vectors were flexed for both the pieces. With continued swaging, the material flow changed its character, i.e., it rotated backwards in the neutral zone, due to the effect of the already swaged deformation strengthened material volume. For the cold-swaged piece, the neutral zone was located in the reduction zone (Figure 6a), while for the warm-swaged piece, the position of the neutral zone shifted towards the feeding zone of the swaging dies (Figure 6b).

**Figure 6.** Vectors of material flow for the cold-swaged piece (**a**) and the warm-swaged piece (**b**).

The finally evaluated predicted parameters were stress-related ones, stress intensity during swaging, and residual stress. Figure 7a,b depicts stress intensity via Von Mises stress throughout the axial longitudinal cut for the cold and warm swaged piece, respectively. As can be seen in Figure 7a, the stress intensity during cold-swaging increased to 750 MPa, while for the warm-swaged piece, the maximum stress intensity reached approximately 350 MPa.

**Figure 7.** Stress intensity in axial longitudinal cut depicted via Von Mises stress for the cold-swaged piece (**a**) and the warm-swaged piece (**b**).

Stress can be present in the material not only during the actual processing, but residual stress can also be present after the material left the swaging dies' reduction zone. Figure 8a,b depicts the predicted distribution of residual stress across the cross-section of the cold and warm swaged piece, respectively. Mutual comparison of the figures reveals similar distribution of residual stress within both of the swaged pieces, its intensity corresponding to the intensity of the imposed effective strain for both pieces. Deeper insight into the residual stress distribution was subsequently provided via experimental analysis.

**Figure 8.** Predicted residual stress intensity across cross-section of the cold-swaged piece (**a**) and the warm-swaged piece (**b**).

#### *3.2. Structure Analyses*

Figure 9a–c shows the SEM-BSE scans of the tungsten pseudo-alloy after sintering, cold-swaging, and warm-swaging, respectively. As can be seen, the sintered state exhibited individual round-shaped W agglomerates surrounded by the NiCo matrix (Figure 9a). The structures after cold and warm swaging (Figure 9b,c) exhibited visible differences in the shapes of the agglomerates. While the grains of both the matrix and tungsten deformed significantly during cold swaging due to the substantial imposed shear strain, the elevated temperature during warm swaging decreased the matrix flow stress, which caused the softer matrix to preferentially consume the imposed strain. For this reason, the tungsten agglomerates remained more or less of round shapes.

**Figure 9.** SEM-BSE structure scans for sintered state (**a**); cold-swaged piece (**b**); warm-swaged piece (**c**).

Figure 10a,b shows residual stress depicted via internal misorientations from 0◦ to 15◦ for the cold and warm swaged pieces, respectively. For both pieces, the presence of residual stress was detected in the peripheral regions of the tungsten agglomerates. This phenomenon can be attributed primarily to substructure development and accumulations of dislocations, as proven by the further discussed TEM structure characterisation.

**Figure 10.** Integral grain misorientations for ChemEngineering-610706 cold-swaged piece (**a**) and the warm-swaged piece (**b**).

As shown by the computational analyses of material flow, the increased temperature of 900 ◦C imparted a more intense plastic flow, which imparts a higher probability of fixation of elastic strain within the warm-swaged material. In other words, the warm-swaged piece (Figure 10b) exhibited a higher volume of locations featuring high internal misorientations, i.e., internal stress, than the cold-swaged piece (Figure 10a). As can be seen from Figure 9b, the misorientations occurred not only at the peripheries of the agglomerates, but also within the more ductile NiCo matrix.

Figure 11a,b depicts TEM scans of the cold- and warm-swaged pieces, respectively. As can be seen, the cold-swaged piece exhibited distinctive interfaces between the tungsten agglomerates (dark area) and the NiCo matrix (light area). On the other hand, the warm-swaged piece exhibited diffusion of tungsten to the NiCo matrix; the areas within the matrix with increased tungsten content can be seen in the figure as "dark shadows" within the light matrix area. To confirm the presence of tungsten within the matrix, a line scan measuring chemical composition through such a darker shadow within the matrix was carried out. Figure 11c depicts the area in which the line scan was carried out (depicted by the straight line). Figure 11d then shows the results of the line scan as regards the chemical composition of the investigated area. As is evident, the size of the interface between the agglomerate and the matrix after warm swaging and the content of tungsten within the matrix in the vicinity of the agglomerates reached up to 50 wt.%.

**Figure 11.** *Cont*.

**Figure 11.** TEM scan showing agglomerate/matrix interface for the cold-swaged piece (**a**) and the warm-swaged piece (**b**); measured line scan area within the warm-swaged piece (**c**); and chemical composition of investigated line scan (**d**).

#### *3.3. Mechanical Properties*

The mechanical properties of both the swaged pieces were evaluated via tensile tests; the ultimate tensile strength (UTS) and elongation till failure for all the investigated pieces are summarized in Table 1.

**Table 1.** Mechanical properties resulting from tensile tests.

As can be seen, the powder-based sintered material exhibited the lowest recorded UTS of approximately 860 MPa; however, the elongation till failure reached up to ~20%. The highest strength was recorded for the cold-swaged piece, which it exhibited a UTS of 1828 MPa after the second swaging pass. The UTS after warm swaging increased too, but the UTS value was slightly lower than after cold swaging (1650 MPa vs. 1828 MPa). The fact that the increase in UTS occurred at the expense of plasticity for both the swaged pieces needs to be stressed. In other words, the decrease in plastic properties after swaging was evident in both cases (especially with higher imposed strain). However, the more favourable mutual combination of strength and plastic properties was achieved after warm swaging.

#### **4. Discussion**

The numerical simulations revealed certain differences in the material behaviours of both the work pieces during swaging. For successful swaging of the cold-swaged piece, no preheating was necessary, which makes the swaging process more effective from the viewpoint of processing time. On the other hand, induction heating of the warm-swaged piece is considerably quicker when compared to heating in a furnace and homogeneous heating of the work piece to the swaging temperature only took 24 s. Nevertheless, certain temperature gradient between the surface and axial regions of the work piece developed during heating. The gradient was primarily dependent on the heating time; it was most significant at the beginning of heating and diminished by the time heating was finished.

Among the advantages of induction heating is also that it can favourably be placed directly in front of the swaging machine, which, however, directly affects the material characteristics [33,34], especially the complex material flow. Continuous heating keeps the unswagged work piece in front of the swaging head at the swaging temperature of 900 ◦C, which introduces gradients in mechanical

properties, especially flow stress, between the swaged and unswaged material volumes. As shown by the predicted plastic flow vectors, gradients in temperature-related material properties influenced significantly localization of the neutral zone, the plastic flow vectors in which changed their orientations.

Regardless the effect of induction heating, the predictions revealed temperature variations during processing for both the swaged pieces. The cold-swaged piece exhibited gradually increasing difference in temperatures in the surface and axial swaged piece regions (up to almost 200 ◦C). Such behaviour is typical for cold swaging technology; the surface temperature of the swaged piece increases primarily due to the effect of direct contact of the surface with the swaging dies, i.e., friction effect [35]. Nevertheless, this effect is reduced by the heat transfer through the cold swaging dies. Warm swaging also introduced a certain increase in temperature, primarily in the work piece surface region, where the maximum temperature increased to 990 ◦C. In addition to the effect of friction, the temperature increase during swaging can also be attributed to deformation heat generated by the influence of the imposed shear strain [36].

The material characteristics during swaging are primarily influenced by the effect of the swaging force, featuring two main components: tangential and axial [29]. Whereas the axial swaging force component contributes to the gradual elongation of the swaged piece and provides the swaged product with its final shape, the tangential swaging force component imparts intensive shear strain, resulting in structure changes, introducing an increase in the actual processing temperature, imparting flexure of the plastic flow vectors, but also imparting the gradient of the imposed effective strain from the swaged piece surface region towards its axial region; the effective strain is the highest in the surface swaged piece region, the effect of the rotating dies introducing shear strain, which is the most intense. This phenomenon was evident especially for the cold-swaged piece, the flow stress for which was aggravated. During warm swaging, the flow stress of the THA decreased, reducing the effective strain gradient across the swaged piece cross-section, but also resulted in the higher total imposed effective strain.

The flow stress, together with the differences in the plastic flows of both the swaged pieces, also non-negligibly influenced the stress intensity, and consequently the load on the swaging machine. The predicted stress intensity, pointing to the resistance of the swaged material against deformation, was significantly higher for the cold-swaged piece. The predicted distribution of residual stress was more or less comparable for the cold and warm swaged piece. Nevertheless, the experimental observations showed micro-differences in the distribution of residual stress within the NiCo matrix. The results showed that both the swaged pieces exhibited the presence of residual stress in the peripheral areas of tungsten agglomerates. As proven by the transmission electron microscopy analyses, this phenomenon could primarily be attributed to substructure development within the agglomerates. However, the more intense plastic flow supported by the elevated swaging temperature introduced the presence of residual stress, i.e., unrelaxed elastic deformation, within the matrix of the warm-swaged piece. The primary reason for such behaviour lay in the intense plastic flow imparted by the increased swaging temperature. The harder tungsten agglomerates featuring lower formability acted as "transferring agents" of the imposed strain during swaging. During processing, mutual distances between the individual agglomerates decrease, i.e., the volume of the ductile matrix between the hard agglomerates gradually decreases, by the effect of which the interactions of the individual agglomerates become more intense. Moreover, the occurring diffusion of tungsten to the NiCo matrix, which intensifies with increasing processing temperature, should also be considered. By the effect of this phenomenon, the matrix hardens, and residual stress develops, especially at the W/NiCo interfaces. In other words, the matrix then involves locations featuring varying chemical compositions, imparting local differences within the matrix, and consequently introducing strain gradients resulting in the presence of residual stress.

#### **5. Conclusions**

The study focused on the numerical prediction of the deformation behaviour of a 93W6Ni1Co tungsten heavy alloy rotary swaged at 20 ◦C and 900 ◦C. The study was supplemented with numerical investigations of the applied induction heating, and experimental observations of residual stress and substructure development within the swaged pieces.

The predictions showed the overall time for heating of the THA to 900 ◦C to be 24 s. The induction heating promoted shifting of the neutral zone the plastic flow vectors in which reverse primarily by the effect of flow stress gradient between the swaged and unswagged materials.

Warm swaging at 900 ◦C imparted a more homogeneous distribution of the imposed strain. However, for both the swaged pieces, the highest strain was observed in their surface regions, which were directly affected by the swaging dies. This finding corresponded to the predicted distribution of temperature throughout the swaged pieces; for both, the temperature increased in their surface regions primarily due to the development of deformation heat.

Both the swaged pieces exhibited the presence of residual stress in the peripheral areas of tungsten agglomerates, which could primarily be attributed to substructure development. However, the supported plastic flow of the warm-swaged piece also introduced the presence of residual stress, i.e., unrelaxed elastic deformation, within the warm-swaged piece matrix, and diffusion of W to the NiCo matrix. The UTS was the highest for the cold-swaged piece (1828 MPa), and slightly lower for the warm-swaged piece (1650 MPa); this phenomenon can primarily by attributed to the occurrence of softening of the matrix.

**Author Contributions:** The contributions of the individual authors were the following: methodology, R.K., L.K. (Lenka Kunˇcická), and A.M.; numerical prediction and evaluation, R.K., R.P., A.M., L.K. (Ludmila Krátká); experimental validation, R.K., L.K. (Lenka Kunˇcická), L.K. (Ludmila Krátká), and A.M.; microscopy investigation and evaluation, L.K. (Lenka Kunˇcická); writing—original draft preparation, L.K. (Ludmila Krátká), R.K., L.K. (Lenka Kunˇcická); writing—review and editing, L.K. (Lenka Kunˇcická), R.K.; project administration, A.M.; funding acquisition, R.K., R.P., and A.M.

**Funding:** This work was supported by Ministry of Education, Youth and Sports, Czech Republic in the framework of the projects SP 2019/74 and SP 2019/43 and by the project no. 19-15479S by the Grant Agency of the Czech Republic.

**Acknowledgments:** The experiments were performed using equipment available at the Technical University of Ostrava. We also acknowledge the help of IPMinfra research infrastructure and working team of Institute of Physics of Materials, CAS.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Correlating Microstrain and Activated Slip Systems with Mechanical Properties within Rotary Swaged WNiCo Pseudoalloy**

#### **Pavel Strunz 1,\*, Lenka Kunˇcická 2, Pˇremysl Beran 1,3, Radim Kocich <sup>2</sup> and Charles Hervoches <sup>1</sup>**

Received: 27 September 2019; Accepted: 30 December 2019; Published: 3 January 2020

**Abstract:** Due to their superb mechanical properties and high specific mass, tungsten heavy alloys are used in demanding applications, such as kinetic penetrators, gyroscope rotors, or radiation shielding. However, their structure, consisting of hard tungsten particles embedded in a soft matrix, makes the deformation processing a challenging task. This study focused on the characterization of deformation behavior during thermomechanical processing of a WNiCo tungsten heavy alloy (THA) via the method of rotary swaging at various temperatures. Emphasis is given to microstrain development and determination of the activated slip systems and dislocation density via neutron diffraction. The analyses showed that the grains of the NiCo2W matrix refined significantly after the deformation treatments. The microstrain was higher in the cold swaged sample (44.2 <sup>×</sup> 10−4). Both the samples swaged at 20 ◦C and 900 ◦C exhibited the activation of edge dislocations with <111> {110} or <110> {111} slip systems, and/or screw dislocations with <110> slip system in the NiCo2W matrix. Dislocation densities were determined and the results were correlated with the final mechanical properties of the swaged bars.

**Keywords:** tungsten; rotary swaging; neutron diffraction; dislocations; microstrain

#### **1. Introduction**

Given their excellent mechanical and physical properties, tungsten heavy alloys (THAs) are popular for demanding applications in the military, radiation shields, and highly demanding components such as aircraft counter-balances and gyroscope rotors. THAs usually consist of 90–97 wt.% of tungsten plus other elements, such as Co, Ni, Fe, and Cu [1]. THAs are generally two-phase composites consisting of spherical tungsten particles/agglomerates surrounded by a ductile matrix [2]. Nevertheless, the structure characteristics consequently impacting the performance of the final product can non-negligibly be altered by even the slightest modifications to the processing technology. By this reason, characterization of the occurring structural phenomena, such as the possible presence of adiabatic shear bands (ASBs), determination of the presence of microstrain and residual stress, and characterization of dislocations and active slip systems are of the utmost importance. The first mentioned phenomenon has been investigated quite thoroughly [3–5], but detailed works characterizing other structure phenomena during deformation processing of THAs are scarce [6,7].

Recently, several pieces of research have shown that imposing intensive shear strain into THAs during their production enhances their utility properties and ballistic performance [8,9]. Effective structure refinement can preferably be introduced via methods of severe plastic deformation (SPD), for example via the widely researched equal channel angular pressing (ECAP) and its modifications, which have been proven to effectively refine the grain size down to several hundreds of nanometres for various materials from aluminium to tungsten [10–16]. Nevertheless, the major drawback of prospective industrial use of SPD methods is their limited applicability for bulk volume samples. For example, the most effective high pressure torsion (HPT) method is only suitable to process coin-like samples [17]. By these reasons, THAs are mostly fabricated via various thermomechanical treatments and more conventional technologies, such as hot extrusion [18,19], cold rolling [20], and swaging [21–23].

Rotary swaging (RS) is an intensive plastic deformation method advantageously used in the industry to gradually reduce cross-sections and increase lengths of axisymmetric workpieces [24,25]. Given by its incremental character and dominant compressive stress state, the method can favourably be used to process sintered materials [26]. The dominant shear strain mechanism enables elimination of residual porosity and imparts significant structure refinement.

The primary aim of the presented study was to determine microstrain and characterize the dislocations and active slip systems in the original sintered THA, as well as in the rotary swaged bars, in order to characterize the effects of thermomechanical treatment on the microstructure and final mechanical properties. Throughout the paper, the term "microstrain" denotes the root mean square of the variations in the lattice parameters across the individual crystallites across microscopic distances (root mean square strain, RMSS). In contrast, the term "macrostrain" (not reported here) refers to the overall change in the lattice parameter caused, for example, by a residual stress distribution across the whole component. Generally, microstrain can be caused by a distribution of crystal defects such as vacancies, dislocations, stacking or twin faults.

Neutron powder diffraction method was used as the principal tool for the characterization of structure and microstructure. The strength of neutron diffraction lies in the possibility to provide information from the bulk of the sample, not only from its near-surface region. This fact is very important, especially for THA, where the material is composed mainly of tungsten which is highly absorbing other radiation (X-ray, electrons). When neutrons are used, the signal is averaged over a large volume and the effects of local variability, large grain size and possible local artefacts are minimized.

#### **2. Materials and Methods**

The W-Ni-Co (93-6-1) wt.% (80.9-16.4-2.7 at.%) pseudo-alloy was produced by powder metallurgy. The particle size of initial W, Ni and Co powders was in the range of 2–4 μm. The weighted mixture of the powders was homogeneously mixed and then sintered at 1500 ◦C under H2 protective atmosphere, and subsequently quenched in water. The as-sintered material, i.e., bars with approximately <sup>12</sup> <sup>×</sup> 18 mm2 elliptical cross sections, is denoted as W\_0 throughout the following text.

The sintered bars were further processed by rotary swaging (RS) into circular swaged bars with a diameter of 10 mm. RS was performed in two different ways: at room temperature (sample W\_A) and at 900 ◦C (sample W\_B).

The neutron diffraction patterns for structure and microstrain determination were collected at ambient temperature on the MEREDIT diffractometer of CANAM infrastructure at NPI Rež near ˇ Prague [27]. A mosaic Cu monochromator (reflection 220) provided neutrons with a wavelength of λ = 1.46 Å. A small (0.4%) λ/2 (0.73 Å) contamination of the incoming beam was present and was taken into account during the analysis. The samples were fixed in the beam using a sample holder enabling a sample rotation along the vertical axis to average the texture and large-grain influence on the diffracted intensities within the diffraction plane. A neutron beam size was selected to submerge the sample fully in the beam. The diffraction patterns were collected from 4 to 144◦ of 2θ (where θ is the scattering angle) with a step size of 0.08◦ using a multi-detector bank (35 3He point counters with corresponding 10' Soller collimators). In order to assess the grain size, an additional neutron diffraction measurement was performed using the TKSN-400 diffractometer [28] equipped with a 2D position-sensitive detector. This measurement was carried out with the neutron wavelength of λ = 1.21 Å.

Further sample analyses were performed using scanning and transmission electron microscopy (SEM and TEM) on ion-polished transversal samples taken from both the swaged pieces and from the original sintered material. SEM-EBSD (electron backscatter diffraction) analyses were performed using a TESCAN Lyra 3 device equipped with NordlysNano EBSD detector with the scan step of 0.1 μm. The substructures and grains analyses were performed using ATEX [29] and Channel 5 software. TEM images were acquired on ion-polished thin foils with a JEOL 2100F device.

The last step was characterization of mechanical properties via tensile tests performed to evaluate the mechanical behavior of the sintered and swaged material states and to determine their ultimate tensile strength (UTS) and maximum elongation. Tensile testing was performed with 100 mm long bars and a strain rate of 1.3 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> using a Zwick device. By the reason that determination of elastic properties on tungsten heavy alloys is complicated by tensile tests due to possible deflections of the stress-strain curves, the elastic moduli were additionally determined via ultrasound measurements by an Olympus 38DL Plus device which applies the Pulse Echo Overlap (PEO).

#### **3. Results**

#### *3.1. Phase Identification and Grain Size*

Prior to microstrain determination, phase identification was done using diffractograms measured at neutron diffractometer MEREDIT [27] for all the samples. An example of the measured and calculated neutron diffraction pattern for the sample W\_B with recognized phases is shown in Figure 1. The other two samples exhibited similar phase composition and the diffractograms are similar, although they differ in details due to peak broadening, as will be discussed later in the text. The phase identification and analysis in all the samples was performed by full-pattern refinement using FullProf software [30].

**Figure 1.** Measured and calculated neutron diffraction pattern for W\_B sample used for phase determination as well as for microstrain characterization. The Bragg positions for individual recognized phases (from the top W-B2 and NiCo2W; it should be noted that also the peak positions for λ/2 contamination wavelength are shown) are below the intensity curve. The difference between the measured and calculated intensity is shown as well.

Two phases were identified in all samples. The main phase is α-W (B2 structure; within the text, it is referred to as W-B2 phase). The second phase with a weight fraction of 6%–7% has pure-Ni-like structure (fcc) with the lattice parameter of about 3.60 Å. The lattice parameter of this Ni-like phase is slightly larger than the one for pure nickel (3.55 Å), thus indicating alloying with larger W atoms

(atomic radius of W, Ni and Co is 139, 124 and 125 pm, respectively). During the structural refinement, it was assumed that the second phase consists of Ni and Co in the same ratio as the initial composition (6:1) with a further addition of 2 at.% of W. This secondary phase is denoted NiCo2W in what follows.

It was found using data from 2D detector of TKSN-400 diffractometer [28] that the W\_0 bar (i.e., the sample without rotary swaging forming) has a fine-grained microstructure for the W-B2 phase but very raw-grained microstructure of the NiCo2W phase. The NiCo2W phase produces spots on the 2D detector while the W grains of the W-B2 phase result in a smooth Debye–Scherrer diffraction conus, as can be seen in Figure 2a. Taking into account the gauge volume of 0.13 cm3, detector characteristics, and the geometrical arrangement of the experiment, the grain size of the NiCo2W phase in the W\_0 sample can be estimated to be in the range 0.2–1 mm.

**Figure 2.** Part of the diffractogram taken with 2D detector at TKSN-400 diffractometer (in the angular range 12.5–22◦). (**a**) W\_0 sample: Left strip, W-B2 110 reflection; right spot, NiCo2W 111 large-grain reflection; (**b**) W\_B sample: The smooth strips of intensities from W-B2 110 (left) and NiCo2W 111 (right) reflections of fine-grained phases after rotary swaging.

The large-grain microstructure is refined by rotary swaging. W\_A and W\_B samples already exhibit the fine-grained NiCo2W phase, as can be seen in Figure 2b, taken for the same angular range as Figure 2a. The spotty pattern of the NiCo2W phase changed here to a smooth pattern of fine-grained NiCo2W 111 reflection on the right side of the angular range while the character of the W-B2 110 reflection (on the left) remained unchanged after rotary swaging.

The finding of the large-grain microstructure of the NiCo2W phase in the W\_0 sample bar stressed the necessity to rotate the samples around the vertical axis in order to minimize the influence of the large grains on the resulting diffractogram and consequently on the microstrain determination.

The results of neutron diffraction analyses were further supported via electron backscattering observations. Figure 3a,b shows EBSD scans depicting the orientations of the NiCo2W grains in W\_A and W\_B samples, respectively. The depicted colours in the unit triangle indicate the orientation of the axis normal to the investigated sample surface in the crystal reference frame. With respect to the W\_0 sample (not shown here), the size of NiCo2W grains was significantly refined and substructure developed. The average NiCo2W grain sizes were 1.35 μm for the W\_A sample and 1.0 μm for the W\_B sample. Regarding the W-B2 grains, the original W powder agglomerated during sintering and formed particles with the sizes of several dozens of micrometres, as can be seen in Figure 3 and as was also reported previously [6].

It comes from the concurrent texture investigation by neutron scattering, which will be published elsewhere, that there was no preferential orientation of the NiCo2W phase in the W\_0 bar. On the other hand, the NiCo2W phase was textured with <111> crystallographic directions preferentially oriented along the sample bar axis after rotary swaging. The texture was significantly stronger for the cold swaged sample W\_A than for the warm swaged sample (W\_B). There is a certain relationship

between the texture and mechanical properties. Nevertheless, the relationship is not described here purposely, as it will be a topic of a detailed study published in a future paper.

**Figure 3.** Electron backscatter diffraction (EBSD) scan of the surface perpendicular to the sample-bar axis depicting NiCo2W phase for sample: (**a**) W\_A; (**b**) W\_B. The depicted colors in the unit triangle indicate the orientation of the axis normal to the investigated sample surface in the crystal reference frame. The gray areas are W-B2 agglomerates.

#### *3.2. Data for Microstrain Determination*

Figure 4a–c display the zoomed selected angular range of the measured and calculated diffractograms for all three samples with indexed reflections for both W-B2 and NiCo2W phases. In order to display the extent of sample peak broadening, Figure 4d additionally shows W\_B sample data together with a pattern calculated without any sample broadening effect.

**Figure 4.** Detail of measured and calculated diffractograms showing the same indexed reflections from both W-B2 and NiCo2W phases in (**a**) W\_0, (**b**) W\_A and (**c**) W\_B samples. (**d**) Additional W\_B sample data together with theoretical simulation of a sample profile without any microstrain present in the NiCo2W phase (i.e., only the instrumental broadening effect is present).

By comparing the measured peaks (see Figure 4) of the W-B2 phase for the sample without deformation and samples after the RS process, no significant changes are recognized. On the other hand, NiCo2W reflections broadened significantly with respect to pure instrumental broadening (see Figure 4d). This indicates an increase in microstrain and dislocation density in NiCo2W phase while the W-B2 phase was not affected significantly.

The microstructural fit to the measured data is discussed in Section 4.

#### *3.3. Electron Microscopy*

Figure 5 shows a TEM scan of the W\_A sample taken in a location near the NiCo2W/W-B2 interface. The substantial presence of dislocations within the NiCo2W matrix was observed.

**Figure 5.** Transmission electron microscope image of W\_A sample NiCo2W phase.

#### *3.4. Material Properties*

The stress-strain curves for W\_0, W\_A, and W\_B samples are depicted in Figure 6. The sintered W\_0 sample exhibited the lowest UTS of approximately 860 MPa. On the other hand, the sample featured a relatively high maximum elongation of more than 18%. As a result of the intensive imposed shear strain, the strength increased substantially; however, plasticity (maximum elongation) decreased after both swaging regimes. The total strengthening, i.e., UTS, was higher for the W\_A sample, whereas the W\_B sample featured higher plasticity.

**Figure 6.** Experimental stress-strain curves for W\_0, W\_A, and W\_B samples.

The physical properties measured via ultrasound are depicted in Table 1 (the average value from 5 independent measurements taken per sample).

**Table 1.** Physical properties resulting from ultrasound measurements.

#### **4. Data Analysis and Discussion**

Detailed analysis of the measured neutron diffraction data was carried out with the intention to determine microstrain, dislocation type and dislocation density. With the support of the data obtained by other techniques, these microstructural parameters are related to the determined mechanical properties.

#### *4.1. Microstrain Determination*

First, a phenomenological approach was used to determine microstrain and the line integral breaths of all the measured reflections. FullProf software [30] enables the fitting of a phenomenological model of peak broadening caused by microstructural features, particularly microstrain and grain size, provided that the instrumental broadening is known across the whole measured 2θ range. The instrument profile dependency on scattering angle was obtained by measuring and fitting the standard SiO2 powder sample in the identical instrument setup. The profile function parameters were extracted and stored in the instrument resolution file and this file was used during structural and microstructural full-pattern refinement. Then, the FullProf refinement results in the sample contribution to the reflection broadening. This sample broadening of the diffraction peaks is reported in the further text.

The reflection profiles of the tungsten grains (W-B2 phase) in the W\_0 sample exhibited no sample broadening. Therefore, no measurable microstrain is present for this phase in the sintered sample. The small broadening of diffraction peaks in the W-B2 phase after rotary swaging was observed. As the W-B2 phase grains were sufficiently large, no grain-size broadening was present. The sample broadening of the W-B2 peaks was then satisfactorily fit using isotropic microstrain. The output was the maximum (upper limit) strain *e* [31] which in fact represents the microstrain as it is connected with the root mean square strain (RMSS, <sup>ε</sup><sup>2</sup> 0 1/2) through a constant scaling factor, <sup>ε</sup><sup>2</sup> 0 1/2 <sup>=</sup> <sup>√</sup> 2/π*e*. The determined W-B2 phase upper limit strain for the W\_A and W\_B samples was *<sup>e</sup>*W\_A <sup>=</sup> 12.4 <sup>×</sup> <sup>10</sup>−<sup>4</sup> and *<sup>e</sup>*W\_B <sup>=</sup> 10.8 <sup>×</sup> <sup>10</sup><sup>−</sup>4, respectively. It can be seen that there is slightly higher microstrain in the W\_A sample than in the W\_B sample.

In the case of the NiCo2W phase, peak broadening was already present in the W\_0 sample. After rotary swaging, the sample broadening effect still significantly increased. Although RS procedure refined the NiCo2W grains, they still remained sufficiently large in the measured samples (see Figure 3). The grain-size broadening was thus not expected. As the reflections from (200) family were visibly more broadened, the anisotropic strain broadening using Stephens formalism [32,33] was used to determine the microstrain contribution. The fit was successfully carried out assuming only microstrain broadening (i.e., no size broadening). Gaussian profile, as usually done for microstrain broadening, was used to satisfactorily describe the sample broadening effect.

The outputs from the refinement using the anisotropic strain broadening are values of integral breadth for the individual reflections of the NiCo2W phase in reciprocal space, β\* = (β cosθ)/λ, where β is the integral breath in 2θ scale, and also microstrain values for the measured diffraction peaks of the NiCo2W phase in all three samples. The sample integral breadth β\* of each NiCo2W reflection is plotted as a function of reciprocal lattice spacing of the particular reflection *d*\* in Figure 7 (classical Williamson-Hall plot [31]), for W\_0 and W\_B samples.

**Figure 7.** Williamson-Hall plot of integral breadths in reciprocal space for W\_0 and W\_B samples shown with linear fits through the points.

When considering only the strain broadening component as mentioned above, the upper limit strain *e* is connected with the integral breath by the formula β\* = 2 *ed*\*, i.e., its average value can be calculated from the slope of the linear dependence of β\* on *d*\* [31]. The linear fits through the points are shown in Figure 7. The determined average sample microstrain for the individual W\_0, W\_A, and W\_B samples is then *<sup>e</sup>*NiCo\_0 <sup>=</sup> 14.2 <sup>×</sup> <sup>10</sup><sup>−</sup>4, *<sup>e</sup>*NiCo\_A <sup>=</sup> 44.2 <sup>×</sup> <sup>10</sup>−<sup>4</sup> and *<sup>e</sup>*NiCo\_B <sup>=</sup> 41.2 <sup>×</sup> <sup>10</sup><sup>−</sup>4, respectively.

It can be seen that the microstrain in NiCo2W phase very significantly (approximately 3 times) increased after rotary swaging. Further, there is a slightly higher NiCo2W-phase microstrain in the W\_A sample than in the W\_B sample. Most probably, swaging at the temperature of 900 ◦C (W\_B sample) enabled a partial dynamic recrystallization of NiCo2W phase.

The anisotropic character of the NiCo2W phase integral breadths, visible in the Williamson-Hall plot (Figure 7), indicates that microstrain is caused by dislocations. Further, the substantial presence of dislocations in the NiCo2W matrix was confirmed by TEM observations (Figure 5). Therefore, the data (i.e., acquired integral breaths for the NiCo2W phase individual reflections) were further analyzed in order to determine the active slip systems and to estimate the dislocation density after the thermomechanical processing.

#### *4.2. Dislocation and Slip System Type*

The fact that dislocation line broadening is usually anisotropic, i.e., depends on *hkl* reflection, is well known (see [34] and references therein). It is given by the anisotropic characters of the displacement fields of dislocations (line defects), which, moreover, are different for different types of dislocations and slip systems. Then, the anisotropy analysis can be in principle used to characterize the particular types of occurring dislocations and activated slip system [35]. Therefore, we tested this possibility also in the NiCo2W phase for the sintered and rotary swaged samples.

The anisotropy is characterized by the dislocation average contrast factors *C*α−*hkl*, which can be calculated with help of the ANIZC program [36] for various types of dislocations of given characters and slip system α and for all the measured *hkl* reflections. The calculations (i.e., determination of the average contrast factor for the possible types of dislocations) were performed for the measured NiCo2W reflections. To perform the calculations, elastic constants of the material were needed [37]. In order to use as precise values of the elastic constants for the NiCo2W solid solution as possible, the assumption of a linear combination of the individual elastic constants of the original constituents, i.e., Ni and Co, present in the ratio of 6:1, was made. The considered values found in the literature were the following: nickel [38]—C11 = 256.5 GPa, C12 = 151.5 GPa, C44 = 123.9 GPa; β-cobalt (fcc) [39]—C11

= 239.8 GPa, C12 = 163.4 GPa, C44 = 133.4 GPa. Their linear combinations used to estimate the elastic constants of the alloy were then: C11 = 246.9 GPa, C12 = 156.1 GPa, C44 = 125.5 GPa.

After calculation of the average contrast factors *C*<sup>α</sup> *hkl*, a modified Williamson-Hall plot [37], i.e., the dependence of β\* on *d*\**C*1/2, can be drawn. The results for the W\_0, W\_A, and W\_B samples are depicted in Figure 8a–o. As can be seen, the best result (concerning fitting the linear dependence) for the W\_0 sample was acquired for screw dislocations and <111> slip system (Figure 8e; the corresponding modified Williamson-Hall plot is marked by a red frame). All the other tested dislocation types and slip systems were far worse.

On the other hand, the slip system of <111> {110} edge dislocations fits the best to the measured integral breadths of W\_A and W\_B samples (Figure 8g,l, respectively, marked by a red frame). Nevertheless, the edge dislocations with <110> {111} slip system (Figure 8f,k), as well as the screw dislocations with <110> slip system (Figure 8i,n), exhibited very good linear fits for the rotary swaged samples, too. The two remaining tested dislocation types and slip systems (screw dislocations and <111> slip system, and edge dislocations with <111> {211} slip system) did not exhibit a sufficiently good fit for any of the investigated samples.

Obviously, certain microstrain was present within the structure of the NiCo2W phase already after sintering and subsequent cooling. The dislocations produced during the sintering/quenching process were predominantly of a screw type with <111> slip system. After rotary swaging (W\_A and W\_B samples), the microstrain increased significantly (~3 times) and the deformation mechanism changed either to edge dislocations with <111> {110} or <110> {111} slip system (the latter occurs typically in fcc structures [40]), or to screw dislocations with <110> slip system.

**Figure 8.** *Cont.*

**Figure 8.** Modified Williamson-Hall plots considering various dislocation types and slip systems for samples: (**a**–**e**) W\_0; (**f**–**j**) W\_A; and (**k**–**o**) W\_B. The red color framed plots represent the best-fitting model, the orange ones represent sufficiently good fits. These dislocation types and slip systems are considered as the most probable for the NiCo2W phase of the pseudoalloy and are used for a further evaluation (see text).

Figure 9a then summarizes the modified Williamson-Hall plots of integral breadths for the best fitting slip systems for W\_0, W\_A, and W\_B samples (edge dislocations with <111> {110} slip system for W\_A and W\_B samples). It should be stressed that the quality of fitting for the edge dislocations with <110> {111} slip system and screw dislocations with <110> slip system was almost equal (see Figure 8). A combination of edge dislocation with screw dislocation system was tested as well for W\_A and W\_B samples. The result for the combination of edge dislocations with <110> {111} slip system and screw dislocations with <110> slip system (the ratio of influence on integral breadth was assumed to be 50%:50%) is shown in Figure 9b. A very similarly good result (not shown here) was obtained for the 50%:50% combination of the edge dislocation with <1,1,1> {1,1,0} slip system and the screw dislocation with <1,1,0> slip system.

**Figure 9.** Summarized modified Williamson-Hall plots of integral breadths for W\_0, W\_A, and W\_B samples shown with linear fits through the points: the edge dislocation with <1,1,1> {1,1,0} slip system for W\_A and W\_B samples (**a**); 50%:50% combination of edge dislocations with <110> {111} slip system and screw dislocations with <110> slip system (**b**).

#### *4.3. Dislocation Density*

Information on dislocation types, slip systems and also on dislocation densities may be very useful for modelling the material behavior using crystal plasticity [41]. Based on the results of line broadening of the NiCo2W phase, dislocation densities ρ can in principle be estimated using various numerical procedures [34,35,42–46].

Powder diffraction techniques, in general, have limited accuracy in determining dislocation densities. Precise descriptions of the instrument profile function and the size of the instrumental broadening are essential for this purpose. However, even for precisely determined instrument parameters, another important aspect is the source intensity and the peak-to-background ratio, which are usually limiting in the case of neutron diffraction. Thus, only a rather limited accuracy on the absolute scale of the estimated dislocation densities can be achieved using the above-mentioned analysis of the neutron diffraction line broadening for the NiCo2W phase. Nevertheless, such information has an important indicative value when comparing relatively the obtained dislocation densities for the individual samples.

For the estimation of the density, we used one of the approaches presented in [35]. The integral breath of each reflection can be approximated by,

$$\beta\_{hkl}^\* = \frac{2\sqrt{2}\left(\ln P\right)^{\frac{3}{2}}}{4\ln P - \ln(\ln P)} \, d\_{hkl}^\* C\_{hkl}^{\frac{1}{2}} \, b \, \rho^{\frac{1}{2}} \tag{1}$$

where *b* is Burgers vector, and factor *P* is related to the correlation in the dislocation arrangement.

Assuming that the dislocation correlation factor *P* [35] can be reasonably estimated, dislocation density ρ can be determined based on Equation (1). Due to the moderate resolution of the diffraction data, *P* could not be determined directly from the diffraction profile analysis. For further calculations, the value 10 for the *P* factor was used which is a reasonable value for Gaussian profiles [35] employed in our analysis of the strain broadening. The estimated error in absolute value for ρ was −20% (in this case, *P* would be equal to 15) and +50% (in this case, *P* would be equal to 5).

Considering the above-mentioned assumptions, the dislocation densities for the most favorable slip system in NiCo2W for all the investigated samples were calculated and they are listed below:

W\_0: 1.7 <sup>×</sup> 1011 cm/cm<sup>3</sup> (screw dislocations with <sup>&</sup>lt;111<sup>&</sup>gt; slip system);

W\_A: 8.7 <sup>×</sup> 1011 cm/cm<sup>3</sup> (either edge dislocations with <110> {111} slip system or <110> screw dislocations);

W\_B: 7.6 <sup>×</sup> 10<sup>11</sup> cm/cm<sup>3</sup> (either edge dislocations with <110> {111} slip system or <110> screw dislocations).

As explained above, absolute values of dislocation densities are burdened by a large error. Nevertheless, a relative comparison of the obtained values of dislocation densities between the individual samples brings significantly more precise information as the relative error of the dislocation density values is expected to be similar for all the samples. Then, it can be seen that the dislocation densities increased approximately 5 times after rotary swaging, and that the dislocation density is 15% higher for the sample swaged at room temperature than for the sample deformed at 900 ◦C.

#### *4.4. Mechanical Properties*

The measured stress-strain curves (Figure 6) are in accord with the results of the microstructural study. Relatively low UTS and high maximum elongation for the sintered W\_0 sample correspond to the coarse grain size of the NiCo2W phase, which, together with the low dislocation density, did not ensure sufficient strengthening.

The plasticity (maximum elongation) decreased after both swaging regimes. Nevertheless, the intensive imposed shear strain imparted significant accumulation of dislocations in the NiCo2W phase and the strength of the material increased substantially. The accumulated dislocations after RS provide hardening to the NiCo2W phase and thus enable effective transfer of the imposed strain to the W-B2 phase. Although the imposed deformation is predominantly accumulated in the NiCo2W phase, it does not mean that the W-grains are not strained at all. Tungsten grains deformation was already observed in the past [19,20,47]. Also in this study, a small level of microstrain within the W-B2 phase in the rotary swaged samples was detected (see Section 4.1).

The W\_A sample, featuring the highest dislocation density and notable presence of microstrain in the NiCo2W phase, exhibited higher total strengthening (i.e., UTS), whereas the W\_B sample featured higher plasticity. Mutual comparison of the stress-strain curves for the W\_A and W\_B samples also revealed a more gradual strengthening (i.e., smoother curve shape) of the W\_B sample. Swaging at the temperature of 900 ◦C enabled the NiCo2W phase to dynamically recrystallize, which was documented by the very small average grain size of 1.0 μm (Section 3.1) and decreased microstrain values (Section 3.2), and provided it with the ability to consume a greater amount of the imposed energy. This finding is in accordance with the results documented by Kataviˇc et al. [48] who reported the hardening rate of the matrix during intensive shear deformation of tungsten pseudoalloys to be more than two times higher than the hardening rate of W particles up to approximately 15% deformation.

#### **5. Conclusions**

This study focused on characterization of the effects of rotary swaging at various temperatures on a WNiCo tungsten heavy alloy via determination of microstrain and characterization of dislocation types and activated slip systems.

The results showed that the original sintered sample consisted of fine-grained spherical W-B2 type agglomerates surrounded by a coarse-grained NiCo2W matrix. The W-B2 agglomerates did not feature any significant microstrain. However, the NiCo2W matrix exhibited microstrain (magnitude 14.2 <sup>×</sup> <sup>10</sup><sup>−</sup>4) resulting mainly from the presence of screw dislocations with <sup>&</sup>lt;111<sup>&</sup>gt; slip system.

Both cold and warm rotary swaging then imparted grains fragmentation for the NiCo2W matrix and resulted in formation of fine-grained structures within the NiCo2W phase. From results presented elsewhere [49,50], it is also clear that the W-B2 phase is refined by rotary swaging. Further, neutron diffraction revealed that the microstrain increased three times in the NiCo2W phase (44.2 <sup>×</sup> 10−<sup>4</sup> and 41.2 <sup>×</sup> 10−<sup>4</sup> for W\_A and W\_B samples, respectively). On the other hand, a rather low level of microstrain was detected in the tungsten W-B2 phase of the composite after the rotary swaging.

The measured mechanical parameters correspond to the results of the microstructural characterization. The swaged samples exhibited substantial strengthening which was primarily caused by the increase in dislocation density (~5× for the 900 ◦C sample, and even approximately 10% more for the cold swaged one) in the NiCo2W phase. The 20 ◦C swaged bar featured the ultimate tensile strength of almost 1900 MPa. It can be concluded from the neutron diffraction that the dominant deformation mechanisms for both the 20 ◦C and 900 ◦C rotary swaged samples were edge dislocations with <111> {110} or <110> {111} slip system, or screw dislocations with <1,1,0> slip system in NiCo2W phase. A combination of the above-mentioned systems is most probable for the NiCo2W phase as these combinations lead to the best modified Williamson-Hall plot.

**Author Contributions:** Conceptualization, P.S., L.K. and R.K.; Data curation, P.S., L.K., P.B. and R.K.; Formal analysis, P.S. and L.K.; Funding acquisition, P.S. and R.K.; Investigation, P.S., L.K., P.B., R.K. and C.H.; Methodology, P.S., P.B. and C.H.; Project administration, P.S. and R.K.; Resources, P.S., L.K., R.K. and C.H.; Software, P.B.; Supervision, P.S. and R.K.; Validation, P.S., L.K., P.B. and R.K.; Visualization, P.S. and L.K.; Writin—original draft, P.S.; Writing—review & editing, L.K., P.B. and R.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by Czech Science Foundation, grant number 19-15479S. Measurements were carried out at the CANAM infrastructure of the NPI CAS Rež, supported through MŠMT project No. ˇ LM2015056, using the neutron source included in the infrastructure Reactors LVR-15 and LR-0, which is financially supported by the Ministry of Education, Youth and Sports-project LM2015074. In the study, equipment acquired within OPCZ.02.1.01/0.0/0.0/16\_013/0001812 project supported by the European Structural and Investment Funds was used.

**Acknowledgments:** The authors are thankful for the technical support of Pravdomil Toman, Božena Michalcová and Alex Pylypchuk.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Fatigue Life of 7475-T7351 Aluminum After Local Severe Plastic Deformation Caused by Machining**

**Petra Ohnistova 1,\*, Miroslav Piska 1, Martin Petrenec 1, Jiri Dluhos 2, Jana Hornikova <sup>3</sup> and Pavel Sandera <sup>3</sup>**

Received: 7 October 2019; Accepted: 31 October 2019; Published: 3 November 2019

**Abstract:** The fatigue properties of thermo-mechanically treated and machined aluminum alloy 7475-T7351 have been studied. The applied advanced machining strategy induced intensive plastic deformation on the machined surface under defined cutting conditions. Therefore, a detailed study of 3D surface topography was performed. Advanced characterization of the material structure and electron back scattered diffraction mapping of selected chemical phases were performed, as well as energy dispersive X-ray analysis of the surface. Advanced mechanical properties of the material were investigated in situ with a scanning electron microscope that was equipped with a unique tensile fixture. The fatigue results confirmed an evident dispersion of the data, but the mechanism of crack nucleation was established. Fracture surface analysis showed that the cracks nucleated at the brittle secondary particles dispersed in the material matrix. The surface topography of samples that had been machined in wide range of cutting/deformation conditions by milling has not proved to be a decisive factor in terms of the fatigue behavior. The incoherent interface and decohesion between the alumina matrix and the brittle secondary phases proved to significantly affect the ultimate strength of the material. Tool engagement also affected the fatigue resistance of the material.

**Keywords:** crack nucleation; fatigue; plastic deformation; surface topography

#### **1. Introduction**

Aluminum and its alloys are used in a wide range of industrial applications. Duralumins, are employed in the aerospace industry. Due to an ideal combination of low density, high strength, good corrosion resistance, and high resistance to fatigue crack propagation, these special alloys take precedence over other structural materials [1]. Aluminum alloys of the 2000 and 7000 series are widely used for primary and secondary aircraft structures, such as frames, spares, and ribs, where any damage has a crucial impact on safety.

The high strength Al-Zn-Mg-Cu 7475 is an alloy with controlled toughness made in the form of sheets and plates, that has an ideal combination of high strength, good fracture toughness, and resistance to fatigue crack propagation. The 7475 alloy has almost 40% greater fracture toughness than the previous version, 7075 [2]. This progress in mechanical properties is a result of the reduction of the content of iron, silicon, and magnesium, and application of thermo-mechanical and heat treatments which achieve a refined grain size [3]. The 7475 alloy, in the form of plates, is usually available in different tempered conditions such as T651, T7351, and T7651 [4]. However, among modifying the chemical composition of the alloys, the mechanical properties of Al-based alloys can also be enhanced via severe plastic deformation (SPD) methods, including high pressure torsion (HPT) [5], equal channel angular pressing (ECAP) and its modifications [6–8], or rotary swaging (RS) [9].

As the significant portion of components made of alloy 7475-T7351 are processed by different machining strategies, it is difficult to determine the mechanisms of the severe plastic deformation imposed by the each machining process.

The machining process is generally characterized by chip formation as the direct result of the force interactions between the tool and the workpiece. To achieve chip separation, the stress applied between the tool and the workpiece in the chip formation zone must exceed the ultimate strength of the workpiece material. This force interaction is highly dependent on the studied material, the geometry and material of the tool, the cutting environment, and on the defined cutting conditions. However, during chip formation, three zones of plastic deformation can be always observed: (1) a primary deformation zone in front of the cutting tool edge with extreme shear deformation (γ = 2–5) and deformation rates (103–108 s−1), (2) a secondary deformation zone between the chip and the tool rake face, and (3) a tertiary deformation zone between the machined surface and the tool flank face. In the zone of primary plastic deformation, the chip reaches a high temperature over its entire cross-section (in some cases near to the melting point). As a result of this elevated temperature, the metallurgical and mechanical properties change, the chip softens, the frictional force and the cutting resistance between tool and workpiece decreases, the shear plane angle increases, the chip cross section becomes thinner, and the chip speed increases. In the zone of tertiary plastic deformation, the cutting edge of the tool initiates a stress concentration at the contact zone between the tool and the workpiece [10–12]. The machined surface layer is subjected to elastic and plastic deformation at a temperature lower than the temperature of recrystallization. Therefore, material is not melted and any material texture change is permanent and the surface layer is hardened. The imposed force interaction and the thermal load result in residual stress concentration at different depths of the surface layer. These residual stresses can be compressive, increasing the fatigue limit, or tensile, reducing the fatigue limit [13]; they can also alter other mechanical and utility properties [14,15].

As it is well known that nuclei of fatigue cracks are mostly observed on the free surface of the loaded component, all previously described factors have to be taken into consideration while choosing an appropriate machining strategy and machining conditions. Initiation of fatigue cracks on the machined surface is usually closely associated with the combination of severe plastic deformation and the presence of material inclusions or severe plastic deformation and significantly deteriorated surface topography [16].

Ojolo at al. [17] presented results of the four-point bending fatigue testing of a end-milled specimen of alloy 2024. The important fatigue life increase (from 2.67 <sup>×</sup> 103 to 3.6 <sup>×</sup> 103 cycles to failure) was observed for samples machined with use of higher cutting speeds (from 3.77 m/min to 48.25 m/min). A decrease of the surface roughness was observed with an increase of the cutting speed, which may be the result of the thermal softening effect due to accumulated heat which caused a temperature rise in the machining zone. The feed speed was described as the most influential factor affecting the fatigue life. The increase of fatigue life was observed with a decrease of the feed speed (from 60 mm/min to 7 mm/min). Surface topography, on the other hand, deteriorated with increase of the feed speed. Ojolo therefore supposed that while using higher feed speeds, the teeth of the end-mill cutter do not perform perfect swiping of the entire surface of the machined zone to make a perfectly smooth surface. It was also observed that an increase in rake angle from 30◦ to 45◦ resulted in a better surface finish and increased the fatigue life of the specimens (from 2.53 <sup>×</sup> 103 to 3.49 <sup>×</sup> <sup>10</sup><sup>3</sup> cycles to failure).

Many other studies have been devoted to determining the effect of machining strategies on the fatigue life. Some have demonstrated the influence of the surface topography. The effect of machining and surface integrity on the fatigue life has been summarized by Novovic et al. [18]. They reported that a large dispersion of results has been found in the literature; however, fatigue life increase was mostly observed with decreasing surface roughness. Koster [19] found that for a roughness parameter (*Ra*) between 2.5 and 5 μm, the residual stress imposed by the machining process is the most important factor affecting the fatigue life of structural alloys. Koster [19] also reported that this effect was suppressed

by elevated temperature, which allowed relaxation of imposed residual stresses. However, a study of progressive milling technology on the surface topography and fatigue life of aluminum alloy 7475 of Piska et al. [20] showed that in the case of the presence of material inclusions or secondary phases larger than standard topography parameters (such as average roughness of the profile, maximum depth of the valley of the roughness profile, and others), the effect of the surface topography is usually suppressed.

Regarding ever-increasing safety requirements, it is necessary to carefully analyze the effect of the surface quality, including material structure and surface topography together with residual stresses and severe plastic deformation, imposed by the machining process before releasing components into operation.

This study is, therefore, focused on the influence of the different cutting conditions and tool inclination of the face milling strategy applied on the bottom wing panel made of alloy 7475-T7351 with regard to its fatigue life during operational use. The main goal of this study is therefore to define the milling condition range that allows maintaining the optimum balance between the productivity of the production process, the quality of machined surface, and the required fatigue properties.

#### **2. Experimental Materials and Methods**

#### *2.1. Material*

The aluminum alloy 7475-T7351 in the form of 70 mm thick plates was used in this study. The heat treatment with designation T7351 denotes solution heat treatment at 470 ◦C, water quenching, controlled stretching, and artificial ageing (over-aged in two stages: first at 121 ◦C for 25 h, second at 163 ◦C for a period of 24–30 h). The average chemical composition of the alloy is presented in Table 1 and its basic mechanical properties are shown in Table 2.

**Table 1.** Chemical composition of the 7475-T7351 alloy (in wt.%) [4].

**Table 2.** Mechanical properties of the 7475-T7351 alloy [4].

The electron back scattered diffraction mapping (EBSD) study showed a heavily deformed structure with high anisotropy and texture of the grains (and very fine subgrains), as shown in Figure 1.

**Figure 1.** Structure of aluminum alloy 7475-T7351 determined by electron back scattered diffraction mapping (EBSD): (**a**) Longitudinal direction; (**b**) Transversal direction.

During STEM (Scanning Transmission Electron Microscopy) lamella analysis, three different secondary phases have been observed in the material matrix, as indicated in Figure 2. Coarse intermetallic particles Al-Cu-Fe (possibly Al7Cu2Fe) [21,22] and Al-Cr-Fe-Cu-Si in the range from 2 μm up to 20 μm were formed during solidification phase. Precipitated Al-Fe-Si and Al-Mg-Cr dispersoids (possibly Al12Fe3Si; Al12Mg2Cr) were formed by solid state precipitation in the grain boundaries. Third, observed secondary phases can be described as fine metastable precipitates in the material matrix (sizes from 2 nm up to 0.6 μm) and these are responsible for strengthening of the alloy (via *GP*, η´ or η) [22,23].

**Figure 2.** STEM lamella of aluminum alloy 7475-T7351—occurrence of secondary phases.

Energy dispersive X-ray spectroscopy (EDX) was used for elemental analysis of the large intermetallic particles, as shown in Figure 3.

**Figure 3.** EDX of a selected intermetallic particle in alloy 7475-T7351.

#### *2.2. Tool Geometry*

The end-milling whole carbide tool -16×55-115 mm JHF 980 Special provided by SECO Tools company, with (Ti, Al)N coating was used for advanced high feed face milling of the specimens. To exclude any potential impact of inaccurate tool geometry, optical 3D tool geometry analysis was performed using a special software subprogram of the ALICONA-IF G5 optical microscope called "Alicona Edge Master". The principle of this analysis was a gradual positioning of the reference plane perpendicular to the cutting edges of the tool. The measured results of the intersection of the reference plane with the cutting edges were statistically processed to obtain final results of the mean radius of the mean cutting edges and information about the mean cutting angles.

#### *2.3. Surface Topography Analysis*

Complex measurement of the surface topography was performed on a set of samples machined by different cutting parameters of face milling, as presented in Table 3, with the tool positioned perpendicular to the machined surface. Measurement of the surface topography for a set of samples with tool inclination of 1◦ was also performed. The high-resolution optical microscope ALICONA IF-G5 was used for analysis of roughness parameters ("*R*"), waviness parameters ("*W*"), Firestone–Abbott parameters, and other advanced 3D surface texture parameters ("*S*"). The measurement methodology was based on the combination of the small depth of focus of the optical system with vertical scanning. In order to perform complex detection of the surface, the high-precision optics moved vertically along the optical axis and continuously captured data from the surface. A corresponding algorithm converted the acquired sensor data into 3D information and true colour images with full depth of field [24]. Nonmeasured points in the datasets were not taken into consideration for further processing or for calculation of corresponding parameters due their low ratio (flat surfaces of samples, good fits of data with the Gaussian distribution, very low occurrence of nonmeasured points in the whole dataset). The measurement methodology was in accordance with the standard EN ISO 25178-606 [25].

*ap—* axial depth of cut, *ae* — radial depth of cut, *fz* — feed per tooth

#### *2.4. Force Loading Analysis of High Feed Face Milling and Induced Severe Plastic Deformation*

Cutting experiments for various cutting speeds and cutting feeds were carried out with a five-axis milling center MCV 1210/Sinumerik 840D. A stationary KISTLER 957B/SW dynamometer was used for measurement of the force loading during high feed face milling for the different cutting conditions, as presented in Table 4. The results have been analyzed with DynoWare software (type 2825A, Kistler, Wintherthur, Switzerland), where mean values of the maximal instantaneous force loading in the X, Y, and Z directions were used for graphical determination of the resultant force *F1M* and its vector decomposition to the cutting force *FC* and the force perpendicular *FCN*. The cutting force and non-deformed chip cross section *AD* was used for calculation of the specific cutting energy *kc* for a given material characterized by the constants *co,* the axial depth of cut *ap*, the radial with of cut *ae*, a angular tooth engagement φ and the effect of chip thickness on specific force loading expressed with the parameter *mc*:

$$k\_c = \frac{F\_c}{A\_D} = \left(\frac{c\_o}{a\_p \times a\_c}\right) \times \int\_{q\times 1}^{q\times 2} \sin^{1-mc} q \times d\varphi \tag{1}$$

**Table 4.** Set of specimens used for force loading analysis.

The basic model of continuous chip formation and the individual parameters for shear deformation and rate of the deformation can be seen in Figure 4.

**Figure 4.** Analysis of the chip formation model when cutting: (**a**) Continuous chip; (**b**) Position of the measured lamella; (**c**) The shear deformation of the lamella; (**d**) Geometry of the cutting tool; (**e**) Model of the shear deformation in the primary zone; (**f**) Speed vector diagram.

The plastic flow of the material is defined with the condition of a constant volume of machined material V passing through the first deformation zone and converted to the chip:

$$\text{Div}\times V = 0\tag{2}$$

$$V = A\_D \times v\_c = A\_{Dc} \times v\_{ch} \tag{3}$$

where *AD* is the cross section of the undeformed material entering with cutting speed vc and *ADC* is the cross section of the material converted to a chip, leaving with speed of chip vch. The other variables can be understood according to Figure 4a–f.

The angle of the shear plane ϕ is defined as

$$\phi = \operatorname{arctg} \frac{\sin \delta\_o}{\Lambda - \cos \delta\_o} \tag{4}$$

where the orthogonal cutting angle δ*<sup>o</sup>* is sum of the orthogonal flank angle and orthogonal cutting edge angle β*o*, i.e., δ*<sup>o</sup>* = α*<sup>o</sup>* + β*o,* and Λ means the chip thickness coefficient, which is expressed as

$$
\Lambda = \frac{h\_{\rm DC}}{h\_{\rm D}} \tag{5}
$$

The shear deformation γ in the primary zone can be derived as function of shear angle ϕ and orthogonal rake angle γ*<sup>o</sup>*

$$\gamma = \frac{\cos\_{\mathbb{Y}\_\sigma}}{\sin \phi \cos(\phi - \gamma\_\sigma)}\tag{6}$$

and the rate of shear deformation sequentially as

$$\dot{\chi} = \frac{\cos \gamma\_o}{\cos(\phi - \gamma\_o)} \frac{v\_c}{T\_s} \tag{7}$$

The average thickness *Ts* and thickness *hDC* of the material lamella can be measured and calculated statistically by electron microscopy, as seen in Figure 4, and the orthogonal rake angle γ*<sup>o</sup>* can be measured with the Alicona G5 microscope. The parameter *hD* corresponds to the feed per tooth.

#### *2.5. Fatigue Testing and Frature Surface Analysis*

The objective of the fatigue testing was to examine the influence of the defined cutting conditions of the face milling on the fatigue life, as summarized in Table 5. The effect of tool inclination of 1◦ has also been examined.

**Table 5.** Combinations of cutting parameters of high feed face milling.

The geometry of the fatigue specimens was chosen to achieve the best match with final operational use of the bottom wing panel. The main criteria for appropriate specimen geometry are defined as follows:

All specimens were oriented in the L-T direction of the rolled plate. Specimens were machined at the five-axis MCV 1210/Sinumerik 840D milling center (TAJMAC ZPS, share company, Zlin, Czech Republic/Siemens AG, Erlangen, Germany). Specimens were specially protected against bending or torsion during machining, and excess heating was limited by use of CIMSTAR 597 coolant (Cimcool Industrial Products B.V., Vlaardingen, the Netherlands) of 10% volume concentration, 20 bar pressure, and 20 L/min flow rate. All functional areas of the specimens have been protected against scratches, all sharp edges rounded to 0.3 mm radius, and all other stress concentrators have been removed.

Fatigue testing has been performed at special axial testing machine BISS and parameters of the testing are defined as follows:

The source of fatigue crack nucleation was examined with the scanning electron microscope (TESCAN ORSAY HOLDING share company, Brno, Czech Republic) TESCAN MIRA 3 operating in both secondary and backscattered electron mode. The fatigue crack initiation and propagation mechanism and the integrity of the adjacent surfaces were investigated.

#### *2.6. In Situ Testing*

A specimen with special geometry was designed for in situ tensile mechanical testing. Profiles of the specimen were cut by EDM wire cutting (Electrical Discharge Machining) and polished. Flat functional surface areas were face milled with a special high feed end-mill (SECO tool JHF 980 Special, *fz* = 0.05 mm, *vc* = 200 m/min, *ap* = 0.2 mm), and all sharp edges were rounded and polished.

Testing was performed on a special in situ tensile stage MT1000 made by NewTec (10 kN, a tensile stage) and all analyses were carried out with the SEM TESCAN MIRA 3, equipped with the NewTec SoftStrain software (version 1, NEWTEC, Nîmes, France), as shown in Figure 5.

**Figure 5.** (**a**) SEM TESCAN MIRA 3 and in situ tensile stage MT1000 made by NewTec; (**b**) Specimen clamping in the in situ tensile stage.

#### 2.6.1. In Situ Tensile Testing

The main scope of the in situ tensile testing was to observe the crack initiation and propagation mechanism in the 7475-T7351 alloy. Two main analyses were performed: (a) Observation of the entire functional area of the specimen; and (b) detailed observation of selected intermetallic particles located at the free machined surface of the specimen. Engineering strain distribution under tensile loading in selected particle was analyzed by DIC (Digital Image Correlation) in MERCURY real-time tracking software.

#### 2.6.2. In Situ Cyclic Testing

The aim of the in situ tensile cyclic loading was to simulate cyclic loading under the operation mode and to observe crack nucleation and short crack propagation. Detailed observation of twenty selected large intermetallic particles was performed in parallel. Parameters of the fatigue loading were defined as follow:

#### **3. Results**

#### *3.1. Tool Geometry*

The geometry of the tool and profile roughness of the cutting edges were analyzed and the results are presented in Tables 6 and 7. All parameters complied with manufacturer's specifications. The standard deviations were in the range of 3–4% of the mean values.

**Table 6.** SECO end-milling tool -16×55-115 mm JHF 980 Special geometry.

**Table 7.** SECO end-milling tool -16×55-115 mm JHF 980 Special surface profile roughness parameters.

*Ra—*average roughness of the profile, *Rq—*root-mean-square roughness of the profile, *Rz*—mean peak to valley height of the roughness profile, *Rp*—maximum peak height of the roughness profile, *Rv*—maximum valley depth of the roughness profile.

#### *3.2. Surface Topography Analysis*

The effect of the defined cutting conditions and tool inclination on *R*-parameters of the surface topography was evident. For the set of samples machined with tool positioning perpendicular to the machined surface, a digression of the average values of the profile roughness parameter (measured perpendicularly to the cutting speed, along the feed speed and longitudinal axis of the samples) was observed:

On the other hand, the increase of the profile roughness parameters was observed with the increase of the feed speed (increase of the feed per tooth, *fz*, from 0.05 to 0.90 mm). Some examples of the profile roughness values increasing while increasing the feed speed are mentioned below:

• The maximum valley depth of roughness profile (*Rv)* increased from 13.23 to 16.92 μm (an increase of 22%), as indicated in Figure 7.

**Figure 6.** Example of profile roughness parameters (*Ra—*average roughness of the profile, *Rq*—root-mean-square roughness of the profile) for different cutting speed (at feed per tooth *fz* = 0.05 mm). The parameter *R2* reflects the level of statistical correlation.

No linear function was observed between the cutting speed and the waviness parameters, or between the feed speed and the waviness parameters (*R*-squared parameter varied in the range from 0.1 to 0.7 for different waviness parameters), as shown in Figures 8 and 9.

The surface parameters under different cutting parameters of the face milling were examined (with tool positioning perpendicular to the machined surface). The results of the measurement can be summarized as follows:

**Figure 7.** *Cont*.

**Figure 7.** Selected profile roughness parameters for different feed speeds (at cutting speed *vc* = 200 m.min<sup>−</sup>1): (**a**) *Ra* – Average roughness, *Rq* – root-mean-square roughness; (**b**) *Rt* - maximum peak to valley height of roughness profile; (**c**) *Rv* - maximum valley depth of roughness profile.

**Figure 8.** Example of profile waviness parameters (*Wa*—average waviness of the profile, *Wq*—root-mean-square waviness of the profile) for different cutting speeds (at feed per tooth *fz* = 0.05 mm). The parameter *R2* reflects the level of statistical correlation.

**Figure 9.** Example of profile waviness parameters (*Wa—*average waviness of the profile, *Wq*—root-mean-square waviness of the profile) for different feed speeds (at feed per tooth *vc* = 200 m.min<sup>−</sup>1). The parameter *R<sup>2</sup>* reflects the level of statistical correlation.

**Figure 10.** Surface parameters of the selected area for different feed speeds (at cutting speed *vc* = 200 m.min<sup>−</sup>1): (**a**) *Sa* - average height of the selected area, *Sq* - root-mean-square height of the selected area; (**b**) *Sv* - maximum valley depth of the selected area.

The machined surface can by described by the Firestone–Abbott curve, which indicates the percentage of the material of the profile elements at a defined height relative to the evaluation profile length (*R*) or surface area (*S*). This specific surface criterion is characterized by several parameters. The parameters of the core roughness depth (*Rk* and *Sk*) indicate the volume of the material above the core material which can be worn during operational use. The parameters of reduced peak height

(*Rpk* and *Spk*) describe the mean height of peaks above the core material. Furthermore, the parameters of reduced peak height (*Rpk* and *Spk*) express the amount of the material that will be removed during the initial operational wearing process. The parameters of reduced valley height (*Rvk* and *Svk*) describe the mean depth of the valleys below the core material. Therefore, *Rvk* and *Svk* parameters indicate the ability of the machined surface to retain liquids. The parameter *Rmr1* indicates the fraction of the surface which consists of peaks above the core material, and the parameter *Rmr2* indicates the fraction of the surface which will carry the load [28]. An examples of the Firestone–Abbott curve are presented in Figures 11 and 12.

**Figure 11.** Example of Firestone–Abbott curve of the roughness profile (*fz* = 0.05 mm and *vc* = 90 m.min<sup>−</sup>1).

The machined surface under the defined cutting conditions of the face milling showed the following results:

**Figure 13.** Mean height of peaks above the core material for different cutting speeds (at feed per tooth *fz* = 0.05 mm) (**a**) *Rpk* - mean height of peaks above the core material for a profile; (**b**) *Spk* - mean height of peaks above the core material for a selected area.

Analysis of the effect of the tool inclination on the surface parameters can be summarized as follows:

**Figure 14.** Mean height of peaks above the core material for different feed speeds (at cutting speed *vc* = 200 m.min<sup>−</sup>1): (**a**) *Rpk* - mean height of peaks above the core material for a profile; (**b**) *Spk* - mean height of peaks above the core material for a selected area.

**Figure 15.** An example of the surface topography of the machined samples (*fz* = 0.9 mm, *vc* = 200 m.min<sup>−</sup>1, *ap* = 0.2 mm, *ae* = 6.0 mm, dry cutting): (**a**) Without a tool inclination; (**b**) Tool inclination of 1◦; (**c**,**d**) Corresponding autocorrelations and texture aspect ratio values of the surface structures.

#### *3.3. Force Loading Analysis of High Feed Face Milling and Induced Severe Plastic Deformation*

The parameters of primary shear deformation and its rates were very high, as can be seen in Figure 16. Primary shear deformation decreases with increase of the cutting speed as well as with increase of the feed speed (feed per tooth).

(**b**)

**Figure 16.** (**a**) The calculated mean values of the shear deformation and the rate of shear deformation in the primary zone; (**b**) Specific cutting forces for the given cutting conditions.

An explanation of this phenomenon is not easy and other works dealing with dislocations are ongoing. Meanwhile, we tentatively propose that there exists a small region over which there is a sudden proliferation of high angle boundaries in the microstructure of the material as it is deformed into the chip [29].

#### *3.4. Fatigue Testing and Frature Surface Analysis*

The effect of the milling parameters on the fatigue life was striking even if the surface topography after machining was not the key factor affecting the fatigue crack nucleation.

The important decrease of the fatigue life of specimens machined with higher feed rates while keeping the same cutting speed (an increase from feed per tooth *fz* = 0.05 mm to *fz* = 0.90 mm) can be seen in Figure 17. This decrease of the fatigue life may be caused by the severe plastic deformation achieved in the smallest chip cross sections and machined at the highest cutting speeds (*vc* = 200 m.min<sup>−</sup>1).

**Figure 17.** The effect of the feed speed (feed per tooth) on the fatigue life of alloy 7475-T7351—flat un-notched specimen (*vc* = 200 m.min<sup>−</sup>1).

The effect of the higher cutting speed on the fatigue life of the specimen is evident in the case of combination with lower feed speed, as presented in Figure 18. Use of higher feed speed increases the fatigue life for both low and high cycle fatigue modes. However, this effect is suppressed by combination with a high feed cutting strategy.

Slight inclination of the cutting tool (only 1◦) resulted in a 29–64% reduction of total cycles (for specimens machined with the combination of the cutting speed *vc* = 90 m.min-1 and different feed speeds, as indicated in Figure 19).

Regardless of the cutting conditions, the fatigue cracks were always initiated in the large intermetallic particles which occurred in different morphologies (as large particles, elongated particles, or clusters of intermetallic particles). Fatigue cracks were mostly initiated in the intermetallic particles located in the vicinity of the machined surface regardless of the stress level. An example of the fracture surface is presented in Figure 20.

**Figure 18.** The effect of the cutting speed combined low feed cutting on the fatigue life of alloy 7475-T7351—flat un-notched specimen.

**Figure 19.** The effect of the tool inclination on the fatigue life of alloy 7475-T7351—flat un-notched specimen.

**Figure 20.** Fracture surface for a specimen machined with the following parameters: *fz* = 0.90 mm and *vc* = 90m.min<sup>−</sup>1, stress level 250 MPa.

#### *3.5. In Situ Testing*

#### 3.5.1. In Situ Tensile Testing

Crack propagation from large intermetallic particles to the material matrix was observed even before reaching the tensile strength limit (484 MPa), as presented in Figure 21. Large intermetallic particles were the main source of the local stress concentration regardless of the severe plastic deformation caused by the machining process. Cracks propagated at the angle of 45◦ to the direction of the maximum shear stresses.

**Figure 21.** Observation of particle cracking: (**a**) Intermetallic particle before tensile loading; (**b**) Intermetallic particle at yield strength (415–419 MPa); (**c**) Intermetallic particle at tensile strength limit (484 MPa). The loading axis was horizontal.

The evolution of the engineering strain distribution under tensile loading is shown in Figure 22. The average engineering strain at the area of the intermetallic particle at yield strength was in the range of 0.80% to 0.90%, and upon reaching the maximal engineering strain (1.15–1.20%), local crack initiation was observed.

**Figure 22.** DIC analysis of the engineering strain distribution in the intermetallic particle during tensile loading: at the beginning of the tensile loading, at the yield strength, and at the tensile strength limit.

#### 3.5.2. In Situ Cyclic Testing

During in situ cyclic testing, initial local fatigue cracks were observed in some intermetallic particles before the fatigue life (1000 cycles) was reached, regardless of severe plastic deformation induced by the milling process. Local fatigue cracks were initiated in the core of the intermetallic particles, and with the rising number of cycles, the fatigue cracks propagated locally to the boundary of the intermetallic particles and the material matrix. Fatigue testing was interrupted at the fatigue level of 6000 cycles. Up to this number of fatigue cycles, the short fatigue cracks remained inside the intermetallic particles and did not propagate further. DIC analysis confirmed that once the local engineering strain reached values of 1.15–1.20% (at 1000 cycles, in this case), the local fatigue cracks were initiated, as demonstrated in Figure 23.

**Figure 23.** Engineering strain distribution (**a**) at 100 cycles; (**b**) at 500 cycles.

#### **4. Discussion**

The observations from the experimental machining, surface analyses, and fatigue testing confirm similar results as those of Ojolo et al. [17] and Novovic et al. [18].

Surface topography analysis confirmed that the roughness parameters increase with the increase of the feed speed (feed per tooth). The increase of the cutting speed caused a decrease of the surface roughness parameters. This result partially confirms the observation of Ojolo et al. [17].

The increase of the feed speed increases surface topography parameters such as average height of the selected area (*Sa*), root-mean-square height of the selected area (*Sq*), or maximum valley depth of the selected area (*Sv*). The roughness parameters were found to be similar for both strategies (perpendicular and inclined by 1◦); however, a greater increase of roughness parameters was observed while using higher feed speeds for samples machined with the tool inclined by 1◦. The average height of the selected area (*Sa*) showed higher values for samples machined with the tool inclined by 1◦. This parameter was not adequate, considering that maximum valley depth of the selected area (*Sv*) was higher for samples machined by a tool positioned perpendicularly to the machined surface.

The trends of specific cutting force and shear deformation confirm a reduction of plastic deformation with increasing cutting speed, but more intensive deformation with reduction of feed per tooth. In other words, the intensity of plastic deformation is higher for shallow cuts and higher cutting speeds.

The highest fatigue resistance was observed at samples machined with the highest cutting speed (*vc* =200 m.min<sup>−</sup>1) and lowest feed per tooth (*fz* = 0.05 mm). A decrease of the fatigue life of specimens machined with higher feed rates while keeping the same cutting speed was observed (increase from feed per tooth *fz* = 0.05 mm to *fz* = 0.90 mm). This decrease of the fatigue life may be caused by the severe plastic deformation achieved in the smallest chip cross sections and machined at the highest cutting speeds (*vc* = 200 m.min<sup>−</sup>1).

Slight inclination of the cutting tool (1◦) resulted in reduction of the total cycles for specimens machined with cutting speed *vc* = 90 m.min−<sup>1</sup> and different feed speeds. This reduction may be caused by plastic deformation caused by teeth not engaged in the cut, as in the case of face milling perpendicular to the machined surface. This plastic deformation may impose beneficial compressive residual stresses into the machined surface and thus increase the fatigue life.

Therefore, the cutting conditions can affect the material removal rate, but a more serious impact can be seen in terms of the surface quality and the resistance to mechanical loading. The effect of inclusions is very serious, and materials used for dynamic loading should be carefully analyzed not only in view the surface integrity, but also considering the occurrence of the phases, which confirms the results of Piska et al. [20]. The effect of material hardening and thermal softening when cutting should be studied further in terms of the density and arrangement of dislocations, stacking fault energy, and other atomic hardening or softening mechanisms.

#### **5. Conclusions**

The application of very advanced laboratory facilities yielded the following results:

**Author Contributions:** Conceptualization, methodology, writing, investigations in machining, original draft preparation, project administration, funding acquisition: P.O. and M.P. (Miroslav Piska), material analyses M.P. (Martin Petrenec), investigations in the in situ REM measurements: J.D., fatigue investigations: J.H., P.S.

**Funding:** This research was funded by FME BUT Brno, Czech Republic as the "Research of perspective production technologies", FSI-S-19-6014 and by the project CEITEC 2020, LQ1601, Ministry of Education, Youth and Sports of the Czech Republic.

**Acknowledgments:** This research work was supported by the FME BUT Brno, Czech Republic as the "Research of perspective production technologies", FSI-S-19-6014 and by the project CEITEC 2020, LQ1601, Ministry of Education, Youth and Sports of the Czech Republic.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

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