Recrystallization Behavior of Boron Modified ... - Springer Link

ISSN 0031918X, The Physics of Metals and Metallography, 2015, Vol. 116, No. 3, pp. 309–319. © Pleiades Publishing, Ltd., 2015. Original Russian Text © R.A. Gaisin, V.M. Imayev, R.M. Imayev, 2015, published in Fizika Metallov i Metallovedenie, 2015, Vol. 116, No. 3, pp. 325–336.

STRENGTH AND PLASTICITY

Recrystallization Behavior of BoronModified CommercialPurity Titanium during Hot Deformation R. A. Gaisina, V. M. Imayeva, b, and R. M. Imayeva, b a

Institute of Metals Superplasticity Problems, Russian Academy of Sciences, ul. Khalturina 39, Ufa, 450001 Russia b Bashkir State University, ul. Validi 31, Ufa, 450076 Russia email: [email protected] Received March 28, 2014; in final form May 20, 2014

Abstract—A comparative investigation of the recrystallization behavior of ascast commercialpurity (CP) titanium VT10 with addition of 0.2 wt % boron and of unalloyed CP titanium in the hotrolled and ascast states has been performed. The introduction of boron leads to the formation of uniformly distributed short filaments and particles of titanium monoboride (TiB), which stimulates a substantial refinement of the cast structure. The recrystallization behavior of the alloys was studied after deformation by compression at T = 900–600°C using the EBSD analysis. It has been established that the kinetics of the recrystallization of the ascast VT10–0.2B alloy is substantially accelerated compared with that of the unalloyed VT10 alloy due to the presence of borides. The recrystallization behavior of the VT10 alloy with the initial hotrolled struc ture and of the VT10–0.2B alloy with the ascast structure proved to be approximately similar. Based on the investigations performed, recommendations have been made that should facilitate (and reduce the cost of) the production of finegrained and ultrafinegrained semifinished items of CP titanium. Keywords: titanium alloys, titanium monoboride, microstructure, recrystallization DOI: 10.1134/S0031918X15010056

INTRODUCTION Commercialpurity (CP) titanium in the ascast state is usually subjected to hot forging or rolling in order to obtain relatively finegrained rods, slabs, and other semifinished items suitable for both the practi cal application and for subsequent deformation and thermal treatments aimed at obtaining improved mechanical properties. As is known [1], this standard treatment increases the cost of cast titanium by approximately two times. The grain size obtained after standard hot forging/rolling is usually d = 15–50 µm [1–8]. The refinement of the structure compared to the ascast state improves all of the main mechanical characteristics of titanium and increases its technolog ical plasticity (ductility), which creates possibilities to further refine the structure using deformation meth ods. In particular, in recent years, numerous works were devoted to the use of large plastic deformations for the formation of ultrafinegrained (d < 1 µm) structure in titanium. As a rule, either warm or cold (up to cryogenic) treatment is employed to this end, with a large total degree of deformation using various deformation modes (cyclic and monotonic), such as multiple forging, multiple forging with subsequent rolling, equalchannel angular pressing with subse quent rolling, hydrostatic extrusion, etc. [2–10]. With decreasing temperature of deformation, the degree of deformation needed to obtain a completely

recrystallized structure increases substantially because of the change in the mechanisms of deformation and recrystallization. Upon transitioning from hot to warm deformation, the mechanism of discontinuous dynamic recrystallization is replaced by the mecha nisms of fragmentation and continuous dynamic recrystallization; in addition, the mechanism of defor mationinduced twinning is activated, which favors the formation of new recrystallized grains [3–5, 8]. Upon cold deformation, the deficit of thermal activa tion suppresses mechanisms of recrystallization, and the main structureforming mechanisms are deforma tion twinning and fragmentation; with increasing degree of deformation, only fragmentation is active, since with decreasing size of grains/fragments to d ≈ 0.5–1.0 µm, twinning is suppressed [3, 4, 8]. Meanwhile, from a practical viewpoint, the appli cation of warm and cold deformation to titanium using different loading modes and large degrees of deformation is hardly expedient when considering the multiply increasing cost of the thus obtained semifin ished items. In connection with this, decreasing the number of deformationrelated stages and refining the conditions of the deformation treatment needed to produce a homogeneous finegrained or ultrafine grained structure in CP titanium seems to be a topical problem.

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It has been shown in recent works that the doping of titanium alloys with boron to ~0.1 wt % leads to the formation of short filaments and particles of titanium monoboride (TiB), which stimulates a substantial refinement of the initial cast structure. This increases the ductility in the ascast state and facilitates the development of recrystallization during hot deforma tion [11–14]. As for titanium, the effect of small addi tions of boron on the development of recrystallization during hot deformation has not been studied. This work was aimed at investigating the effect of modification with boron on the mechanical properties and recrystallization behavior of CP titanium (alloy VT10) during hot deformation by compression. The amount of boron to be introduced was chosen based on the results of experiments performed using tita nium alloys [11–14], according to which a significant effect of structure refinement is achieved already at a boron concentration of 0.1–0.2 wt %. In this work, we compare the recrystallization behavior of the ascast VT10 alloy modified by boron with the VT10 alloy without boron in hotrolled and ascast states. EXPERIMENTAL As the materials for the investigation, we chose CP titanium (VT10) in the form of an ordinary hotrolled rod 20 mm in diameter, as well as the alloys VT10 and VT10–0.2 wt % B (further, VT10–B) in the ascast state. Ingots of these alloys weighing about 100 g were melted using the argonarc method in a laboratory fur nace (Edmund Bühler). As the initial materials, a rod of the VT10 alloy and a boron powder (99.5%) pro duced at the OAO AVABOR were used. The temperature of the polymorphic transforma tion of the alloys obtained was estimated based on quenching experiments. To do this, samples of small sections were heated to temperatures in the range of 870–920°C in steps of 10 K, held at the chosen tem perature for 15 min, and quenched in water, after which they were aged at Т = 600°C for 1 h. The tem perature of the polymorphic transformation was esti mated metallographically based on the decrease in the amount of the primary α phase. The mechanical tests by compression were per formed at temperatures Т = 20 and 500–900°C at the initial deformation rate ε' = 10–3 s–1. At room temper ature, the deformation was carried out up to failure; at enhanced temperatures, to the equal degrees of defor mation ε = 60%. The samples for compressive tests had dimensions of 5 × 5 × 8 mm. After the tests at enhanced temperatures, the samples were removed from the furnace to the air in a time of no more than 30 s. The tests were performed using a Schenck Trebel RMS100 testing machine. The curves of the depen dences of the true flow stress on the degree of defor mation were constructed with allowance for the uni

form increase in the transverse section of the samples in the course of the tests. Each point was obtained using two samples. The samples deformed at enhanced temperatures were cut in two along the direction of the load applica tion to study the microstructure in the sections obtained. The microstructure studies were carried out using an Olympus GX51 optical microscope and a Mira3 Tescan scanning electron microscope. In the VT10–B and VT10 samples with the initial cast structure after hot deformation, the volume fraction of recrystallized grains was estimated metallographically with allowance for the fine, equiaxed grains. For the VT10 samples in the hotrolled state, the volume fraction of the recrystallized structure has not been estimated. When studying the initial and deformed microstructure obtained after hot deformation, the method of the automated analysis of electron back scatter diffraction patterns (EBSD analysis) with a scanning step of 0.2–1.0 µm was also used. The low angle grain boundaries with misorientations of less than 2° were ignored taking into account the error of the analysis. When estimating the fraction of high angle grain boundaries, the angles with a misorienta tion more than 15° were taken into account. The EBSD analysis was also used to estimate the average size of recrystallized grains. RESULTS AND DISCUSSION Microstructure in the Initial State Figures 1a–1c show the initial microstructures of the ascast alloys VT10 and VT10–B; Fig. 1d dis plays the microstructure of the VT10 alloy in the hot rolled state. The addition of boron leads to the forma tion of uniformly distributed short filaments and par ticles of titanium monoboride (TiB) [15] up to 100 µm long that are located predominantly at the boundaries of colonies and favor a substantial refinement of the cast structure. The size of the colonies and the length of α plates in VT10–B are approximately 10–100 µm against 100–1000 µm in cast VT10. The microstruc ture of the hotrolled VT10 alloy is characterized by a relatively finegrained predominantly equiaxed struc ture with an average grain size d = 20 µm (Fig. 1d). Figure 2 show the results of measurements of the spectra of grainboundary misorientations in the ini tial states of the alloys that were obtained by the EDSD analysis. The spectra of the grainboundary misorientations in the ascast VT10 and VT10–B alloys are similar; the fraction of the highangle (pre dominantly, special) grain boundaries is 68.7 and 65.9% in VT10 and VT10–B, respectively; the other boundaries are lowangle. The spectrum of grainboundary misorientations in the VT10 alloy in the hotrolled state consists by approximately one third of the highangle grain boundaries of predomi nantly random type (which follows from the continu

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Fig. 1. Initial microstructure of the VT10 and VT10–B alloys obtained by scanning electron microscopy: (a, b) EBSD; (c, d) optical microscopy; (a) ascast VT10; (b, c) ascast VT10–B; (d) hotrolled VT10; (a, b) black lines correspond to high angle grain boundaries; white lines, to lowangle boundaries.

ity of the spectrum of highangle grain boundaries, with small peaks in the regions of 30° and 90° corre sponding to special boundaries), and by twothirds of lowangle grain boundaries (Fig. 2c). The high frac tion of lowangle grain boundaries in the hotrolled VT10 alloy appears to be due to the preceding defor mation treatment. If we compare the dimensions of the microstructure components in the hotrolled state in the VT10 alloy and in the ascast VT10–B alloy (Figs. 1c, 1d), we will see that they are more similar than in the ascast VT10 and VT10–B alloys. However, the grainboundary spectra differ substantially. In ascast VT10–B, highangle special grain boundaries are predominant, whereas in the hotrolled VT10 alloy, the lowangle grain bound aries dominate. THE PHYSICS OF METALS AND METALLOGRAPHY

Temperature of the Polymorphic Transformation The temperature of the polymorphic transforma tion was Tβ = 890–900°C for both alloys. Note that, unlike highpurity iodide titanium, CP titanium VT1 0 has no strictly definite temperature of the polymor phic transformation. According to [1, 16], the poly morphic transformation in VT10 occurs in the tem perature range of 885–900°С [1] or even 880–900°C [16]; i.e., it occurs through the twophase region (α + β), the width of which depends on the concen tration of impurities. The absence of the effect of small additions of boron on the temperature of the polymor phic transformation is explained by the extremely low solubility of boron in titanium and was noted earlier for the Ti–6Al–4V alloy [17]. Vol. 116

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Fig. 2. Spectra of grainboundary misorientations obtained by the EBSD method: (a) ascast VT10; (b) ascast VT10–B; (c) hotrolled VT10; HAGB = highangle grain boundaries.

Mechanical Compressive Properties Tests for compression at room temperature have shown close values of the flow stress in VT10 and VT10–B in the ascast state and slightly higher values of the flow stress in VT10 in the hotrolled state (at the degree of deformation of about 20%) (Fig. 3a). The ductility of the ascast VT10–B alloy and of the hotrolled VT10 alloy proved to be higher than the ductility of the ascast VT10 alloy. In the first case, the samples did not fail after compressive deformation to ε = 60%, whereas the samples of the ascast VT10 alloy failed at ε = 36%. The improvement of the deformability in the ascast state upon the modifica tion by boron appears to be caused by the refinement of the initial cast structure in the presence of borides, which was noted earlier in studies of titanium alloys [11]. The slightly higher values of the flow stress obtained for hotrolled VT10 are explained by the relatively small size of grains; the developed substruc ture; and the low content of highangle special grain boundaries, which are weak barriers for the motion of dislocations. In the VT10 and VT10–B alloys in the ascast state there was revealed a large amount of spe cial grain boundaries, which, in spite of a substantially smaller size of grains/plates in VT10–B, provided close values of the flow stress in both alloys. Figures 3b–3f present truestress–strain curves constructed based on the results of compressive tests at enhanced temperatures. With increasing deformation temperature, the flow stresses for the hotrolled VT1 0 alloy and for ascast VT10–B decrease faster than for ascast VT10; therefore, at Т = 700–900°C the curves corresponding to hotrolled VT10 and ascast VT10–B lie below the curves corresponding to as cast VT10. This is due to the smaller size of grains/plates, at the boundaries of which the plastic flow can start. At Т = 500–600°C, the boundaries of the grains/plates are more efficient barriers for the motion of dislocations than at Т = 700–900°C; there fore, the flow stress at the initial stage in the rod of VT10 and in ascast VT10–B at Т = 500–600°C

proves to be slightly higher than in ascast VT10. With increasing degree of deformation, relaxation of stresses in the region of grain boundaries appears to begin; in addition, as will be shown below, in the VT10 rod and in ascast VT10–B, unlike ascast VT10, the dynamic recrystallization develops already at these temperatures. As a result, with an increasing degree of deforma tion, the flow stresses for all states of the alloys become close to one another at Т = 500°C; at Т = 600°C, they become lower in the VT10 rod and ascast VT10–B than in the ascast VT10. The difference in the flow stress at the initial stage between ascast VT10–B and VT10 rod at Т = 500°C appear to be due to a high resistance to plastic flow of the substructure present in the VT10 rod in the initial state, as well as due to the high fraction of special boundaries in VT10–B, which are relatively easily penetrable for dislocations. On the whole, the relationship between the flow stresses in the alloy states under consideration is deter mined by the grain size and barrier properties of the boundaries of grains and plates; by the presence or absence of a substructure; and, beginning with a cer tain degree of deformation, by the rate of the develop ment of dynamic recrystallization and recovery. For the ascast VT10 at all deformation tempera tures there is observed a curve with an extended stage of strengthening and insignificant softening or with only a strengthening. The stage of softening for the as cast VT10 is observed at low temperatures (~500– 600°C), which, taking into account the asymmetric shape of the deformed samples, should be related to the localization of deformation. For the hotrolled VT10 alloy and for VT10–B, the curves demonstrate an insignificant strengthening (for the VT10 rod, at Т = 700–900°C; for VT10–B, at Т = 500°C), or some softening (for the VT10 rod, at Т = 600–700°C; for VT10–B, at Т = 800–900°C). A significant con tribution to the softening in the process of plastic flow in the hotrolled VT10 and in the ascast VT10 in comparison with ascast VT10 comes, as follows from


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Fig. 4. Microstructure (obtained based on the EBSD data) of the central part of samples of the ascast VT10 alloy deformed by compression at various temperatures, and the corresponding spectra of grainboundary misorientations: (a, d) Т = 900°C; (b, e) Т = 800°C; (c, f) Т = 700°C; (a–c) black lines correspond to highangle grain boundaries; white lines, to lowangle bound aries; HAGB = highangle grain boundaries.

the belowgiven results, from the intense development of dynamic recrystallization and also apparently from grainboundary sliding in the region of recrystallized grains. Effect of Hot Deformation on Microstructure The microstructure investigations of deformed samples have revealed a substantially different recrys tallization behavior of the materials. Figure 4 shows the microstructure of the central part of samples deformed at Т = 700–900°C and the corresponding spectra of grainboundary misorientations obtained for the initial ascast VT10 alloy. It can be seen that the hot deformation mainly leads to a change in the geometrical shape of grains. The recrystallization is developed in the form of localized bands after the deformation at T = 900 and 800°C and is barely devel oped at T = 700 and 600°C. The volume fraction of recrystallized grains (with allowance for the entire sec tion of the sample) does not exceed 15–20% after deformation at Т = 900 and 800°C and 5% after defor mation at Т = 700 and 600°C. The average size of recrystallized grains after hot deformation changes

with decreasing temperature nonlinearly; from Т = 900 to 800°C, it increases from d = 11.7 to 17.1 µm, after which, it decreases to d = 1.8 µm at Т = 600°C (Table 1). The fraction of highangle grain boundaries at all deformation temperatures remains low (26.8–37.7%, Table 2). Note the appearance of small peaks near 90° after deformation at Т = 600–800°C (spectrum of misorientations for Т = 600°C is not shown), which correspond to twin boundaries. The deformation twinning in the ascast VT10 appears to develop to some extent even at these high temperatures, which can be explained by the coarse initial size of grains. Thus, the main processes that accompany the deformation of the ascast VT10 alloy are the change in the geometrical shape of grains and the dynamic recovery, whereas the processes of recrystallization are developed weakly and in local regions. The discontin uous dynamic recrystallization based on the migration of initial grain boundaries appear to be ineffective in the case of coarse initial size of α grains/plates that is characteristic of the ascast state of VT10. Figure 5 shows the microstructure of the central part of samples of the ascast VT10–B alloy deformed at Т = 900–700°C and the corresponding


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Table 1. Volume fraction and average size of recrystallized grains in the VT10 alloy (in the ascast and hotrolled states) and in the VT10–B alloy after hot deformation by compression (ε = 60%, ε' = 10–3 s–1) Deformation temperature, °C

Volume fraction of recrystallized structure, %/average size of recrystallized grains, μm VT10, ascast

VT10–B, ascast

VT10, hotrolled (rods)

33/2.1 50/8.8 65/17.3 75/12.5

–/2.4 –/11.5 –/18.6 –/13.2

≤5/1.8 ≤5/9.2 ≤15–20/17.1 ≤15–20/11.7

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Table 2. Fraction of highangle grain boundaries after hot deformation by compression (ε= 60%, ε' = 10–3 s–1) obtained using the EBSD analysis for the central part of deformed samples Fraction of highangle grain boundaries, % Alloy, state VT10, ascast VT10–B, ascast VT10, hotrolled (rods)

initial state

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26.8 72.8 58.1

* The test was performed on the VT10 alloy in the hotrolled state at the same rate and to the same degree of deformation as in this work [3].

spectra of grainboundary misorientations. The refinement of the initial ascast structure due to the presence of borides qualitatively changes the recrystal lization behavior of the material; the volume of recrys tallized grains increases significantly as compared to the preceding case. Thus, according to the metallo graphic estimate, after deformation at Т = 900°C, this increase was about 75%. With decreasing deformation temperature, the volume fraction of recrystallized grains decreases and after deformation at Т = 800, 700, and 600°C is 65, 59, and 33%, respectively. The average size of recrystallized grains changes with decreasing deformation temperature nonlinearly, as in the case of the ascast VT10 alloy; as the temperature decreases from T = 900 to 800°C, it increases from d = 12.5 to d = 17.3 µm, and then decreases to d = 2.1 µm at 600°C (Table 1). The fraction of highangle grain boundaries is maximum at Т = 900°C (72.8%); with decreasing deformation temperature, it decreases to 48.8 and 46.6% at Т = 800 and 700°C, respectively, and increases insignificantly (to 51.8%) at Т = 600°C. Note weak peaks near 30° and 90° after deformation at T = 900 and 600°C (the spectrum of misorientations for Т = 600°C is not given), which correspond to spe cial twin boundaries. In the first case, they most likely appeared immediately after the termination of defor mation; in the second case, they appeared as a result of deformation twinning. On the whole, the data indicate that the main process at Т = 900°C is discontinuous dynamic recrystallization and that it is developed quite efficiently, ensuring a high volume of recrystallized grains and the predominance of highangle grain boundaries. In the course of deformation at Т = 800– 600°C, the main processes are discontinuous dynamic THE PHYSICS OF METALS AND METALLOGRAPHY

recrystallization and recovery. At Т = 600°C, the deformation twinning apparently begins to play a def inite role in the formation of new recrystallized grains; this follows from an insignificant increase in the frac tion of highangle grain boundaries compared to the case of deformation at Т = 700–800°C. The recrystallization behavior of VT10 in the hot rolled state is close to that of VT10–B (Fig. 6). The volume of the recrystallized structure has not been estimated quantitatively; however, it decreased signifi cantly with decreasing deformation temperature. As in the preceding cases, the average size of recrystallized grains changes nonlinearly; as the temperature decreases from Т = 900 to 800°C, d increases from 13.2 to 18.6 µm, then decreases to d = 2.4 µm at Т = 600°C (Table 1). On the whole, the recrystallization behavior (volume and size of recrystallized grains) of VT10 in the hotrolled state agrees with the earlier work [18]. The fraction of highangle grain boundaries in the VT10 alloy in the hotrolled state changes depending on temperature in a complicated manner. As the deformation temperature decreases from 900 to 800 and further to 700°C, the fraction of highangle grain boundaries decreases slightly (from 58.1 to 50.5 and 48.4%, respectively), then increases to 67% at 600°C and again decreases to 55.5% at 500°C (Table 2). The peaks located near 30° and 90° that are observed after deformation at 900°C correspond to special boundaries, which, as in the previous case, are formed immediately after the termination of the deformation. The main processes that accompany the deformation at Т = 900–700°C are the discontinuous dynamic recrystallization and recovery; at Т = 600 and 500°C, Vol. 116

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a significant contribution to the formation of recrys tallized grains appears to come from deformation twinning. This follows from an increase in the fraction of the highangle grain boundaries as compared to the deformation at Т = 800–700°C and from the shape of the σ–ε curves (Figs. 3b, 3c), which is characteristic of dynamic recrystallization. Table 1 contains the results of estimating the vol ume fraction and average size of recrystallized grains; Table 2 shows the data of the EBSD analysis. The results revealed several interesting facts that require special discussion. It can be seen from Table 2 that the fractions of highangle grain boundaries in the central (most strongly recrystallized) part of the samples of the hotrolled VT10 alloy and of the VT10–B alloy change with the deformation temperature in a similar way. Furthermore, taking into account data on the vol ume fraction of recrystallized grains obtained in this work and earlier [18], it can be concluded that the recrystallization behavior of the hotrolled VT10–B alloy and of the ascast VT10 alloy in the range of deformation temperatures of Т = 900–600°C is simi lar despite the difference in their initial microstruc ture. In the ascast VT10, the recrystallization is developed significantly more weakly than in a hot

rolled rod made of VT10 or in the VT10–B alloy, which is caused by the initial coarsegrain structure; therefore, the main processes that accompany defor mation, as was already noted above, are changes in the geometrical shape of grains and the dynamic recovery, which provides a high fraction of lowangle grain boundaries. The enhanced fraction of highangle grain bound aries in VT10–B compared to VT10 in the hot rolled state after deformation at Т = 900°C appears to be due to a substantial difference in the spectra of grain boundaries in the initial state of these alloys. The high fraction of lowangle grain boundaries in the initial VT10 rod and the rapid occurrence of dynamic recovery caused by the high energy of stacking faults in titanium [19] appear to favor the relatively high frac tion of lowangle grain boundaries after deformation, despite the rapid development of dynamic recrystalli zation. In the VT10–B alloy, the formation of high angle grain boundaries was facilitated by the presence of a high fraction of special boundaries in the initial state, which are finished highangle grain boundaries. The experiments performed made it possible to reveal a new interesting effect related to the size of recrystallized grains. The EBSD measurements have


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Fig. 6. Microstructure (obtained based on the EBSD data) of the central part of samples of the hotrolled VT10 alloy deformed by compression at various temperatures, and the corresponding spectra of grainboundary misorientations: (a, d) Т = 900°C; (b, e) Т = 800°C; (c, f) Т = 700°C; (a–c) black lines correspond to highangle grain boundaries; white lines, to lowangle bound aries; HAGB = highangle grain boundaries.

shown that the average grain size after deformation first increases with a decrease in the deformation tem perature from Т = 900 to 800°C and then decreases at Т = 700°C (Table 1, Figs. 4–6). It can be seen that the average size of recrystallized grains is close after defor mation at Т = 900 and 700°C. This effect can be explained by the fact that the temperature Т = 900°C corresponds to the twophase (α + β) field (Tβ = 890– 900°C), which restricted the growth of recrystallized grains in the process of deformation, whereas the tem peratures Т = 800 and 700°C correspond to the com pletely singlephase α region. Therefore, the rate of growth of the recrystallized grains at Т = 800°C proved to be even higher than at Т = 900°C. Taking into account that the maximum volume of recrystallized grains in all alloys was observed after deformation at Т = 900°C, the idea of the implementation of the first deformation treatment of VT10 in the hotrolled state or of the VT10–B alloy in the quasitwophase region (e.g., at a deformation rate of ~10–2 s–1) seems to be quite reasonable. Thus, as was shown by experi ment, the use of small deformations and low forces can produce a homogeneous finegrained microstruc ture with a grain size of about 10 µm (at a greater THE PHYSICS OF METALS AND METALLOGRAPHY

deformation rate, even with a smaller grain size) and with the predominance of highangle grain bound aries. This will improve the ductility of titanium and, in the recrystallized state, it can be subjected to further deformation treatments at reduced temperatures without a risk of fracture, which will ensure the forma tion of an ultrafinegrained microstructure with, probably, minimum labor inputs. This way seems to be more expedient for obtaining bulk ultrafinegrained semifinished items from CP titanium as compared to the methods based on the use of large warm/cold plas tic deformation without previous grain refinement to ~10 µm. The additions to CP titanium of boron (which forms titanium monoboride) in the amount of ~0.1 wt % lead to a substantial refinement of the cast structure and a significant facilitation of the occurrence of the pro cesses of dynamic recrystallization in comparison with cast VT10. Therefore, the additions of boron in the amount of ~0.1 wt % can, apparently, substantially simplify the production of rods and slabs, and reduce the overall cost of the process of fabricating fine and ultrafinegrained semifinished items of CP titanium. Vol. 116

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With decreasing deformation temperature from Т = 700 to 600°C, the fraction of highangle grain boundaries in the hotrolled VT10 and in the VT10–B increases. Most likely, as was noted, this is related to a change in the mechanism of deformation with decreasing temperature. It is well known that, in CP titanium, because of the specific features of its struc ture and related limitedness of the slip systems in it, the deformation twinning is developed in the process of deformation at enhanced temperatures up to 600°C and higher [19, 20]. The deformation twins lead to the fragmentation of the structure, and the twin bound aries are finished highangle boundaries, which facili tates the development of dynamic recrystallization. The twin boundaries in the process of deformation gradually lose their special misorientation because of the incorporation of lattice dislocations into them [21]; therefore, the peaks that correspond to twin boundaries are almost absent in the spectra of the grainboundary misorientations obtained. With a fur ther decrease in the deformation temperature to 500 and 400°C for hotrolled VT10, the lack of thermal activation sharply decelerates the development of the discontinuous dynamic recrystallization, and the frag mentation and continuous dynamic recrystallization become the main mechanisms of the formation of new grains, which is also stimulated by the development of deformation twinning. Note that, as the temperature decreases from 500 to 400°C, the rate of twinning appears to be reduced, which leads to a decrease in the fraction of highangle grain boundaries (Table 2) [3]. The data on the recrystallization behavior of VT10–B and VT10 in the hotrolled state permit us to formulate some recommendations, the fulfillment of which will make it possible to minimize the magni tude of the stored energy of plastic deformation upon the production of finegrained and ultrafinegrained semifinished items from CP titanium as follows: (1) The development of recrystallization processes is substantially facilitated by the addition of boron in the amount of ~0.1 wt %. (2) It is expedient to carry out the first deforma tionbased processing of the hotrolled VT10 or of the ascast VT10–B in the quasitwophase (α + β) field. At a notverylow rate of deformation, a recrystal lized structure can be obtained in this way with an aver age grain size to 10 µm and with the predominance of highangle grain boundaries using relatively small deformations and low forces during the deformation. (3) It is expedient to carry out the subsequent deformationbased conversion of VT10 or VT10–B with the purpose of obtaining ultrafinegrained struc ture at reduced temperatures corresponding to the high activity of deformation twinning, which appar ently favors the development of dynamic recrystalliza tion [3, 4, 18]. The rate of deformation twinning depends substantially on the degree of purity of tita nium [20]. For the CP titanium of grade VT10, the

most suitable temperature of deformation treatment intended for obtaining ultrafinegrained (or almost ultrafinegrained) state is 600–500°C, where defor mation twinning is characterized by the greatest rate. In this case, the lower the deformation temperature, the higher the degree of deformation required to achieve a completely recrystallized and ultrafine grained structure with the predominance of a high angle grain boundaries. After a completely recrystal lized homogeneous ultrafinegrained structure is obtained, a further decrease in the deformation tem perature will be possible at subsequent treatments. A key condition that can provide an efficient grain refinement is the achievement of the predominance of highangle grain boundaries at each stage of deforma tion (at each temperature) [10]. The fulfillment of these recommendations will increase the ductility of titanium at each stage of deformation and, as can be expected, will substantially facilitate the production of finegrained and ultrafine grained semifinished items from CP titanium. CONCLUSIONS (1) The introduction of 0.2 wt % boron into the titanium alloy VT10 leads to the formation of short uniformly distributed filaments of titanium monoboride (TiB), which favor a substantial (tenfold or greater) decrease in the size of colonies and in the length of α plates in the ascast state. At room temper ature, this leads to a noticeable growth of plasticity upon compression as compared to the ascast VT10 alloy. At enhanced temperature, the shape of the flow stress–strain curves of the VT10 alloy changes signif icantly, i.e., the flow stress changes more gradually and, at Т = 600–900°C, a stage of softening appears in the curves. (2) It has been revealed that the most efficient development of dynamic recrystallization, which ensures the maximum recrystallized volume and the average size of recrystallized grains of about 10 µm, is provided by the hot deformation of the ascast VT10–B and hotrolled VT10 alloys in the quasi twophase (α + β) field (at Т = 900°C). (3) The use of the analysis of electron backscatter diffraction patterns (EBSD analysis) revealed the effect of the initial structure of the VT10 alloy on the arising spectrum of grainboundary misorientations, which makes it possible to use this method, along with the direct microstructure observations, to estimate the development of the processes of dynamic recrystalli zation, recovery, and deformation twinning. It has been shown that the recrystallization behavior of the ascast VT10–B alloy and hotrolled VT10 alloy in the temperature range of 600–900°C is similar. (4) Based on the results of the investigation of the recrystallization behavior of the ascast VT10–B alloy and VT10 alloy in the ascast and hotrolled


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states, practical recommendations are suggested. Fol lowing these recommendations should facilitate, i.e., lower the cost, the production of finegrained and ultrafinegrained semifinished items from CP tita nium. It has been established that the small additions of boron accelerate the development of recrystalliza tion processes upon hot deformation. It has been shown that the development of recrystallization is facilitated if the hot deformation is carried out in the quasitwophase (α + β) field. As was shown for the ascast VT10–B and hotrolled VT10 alloys, the size of recrystallized grains achieved in this case proves to be approximately the same as if the deformation were implemented at a temperature that is 200 K and the volume of the recrystallized structure were signifi cantly greater. It is recommended that the subsequent deformation treatments required to obtain ultrafine grained semifinished items from the VT10 alloy be carried out at temperatures that correspond to the high activity of deformation twinning. In this case, it should be taken into account that, the lower the defor mation temperature, the higher the degree of defor mation required to obtain a completely recrystallized ultrafinegrained structure with the predominance of highangle grain boundaries. REFERENCES 1. Physical Metallurgy of Titanium Alloys, Ed. by N. F. Anoshkin (Metallurgiya, Moscow, 1980) [in Rus sian]. 2. J. L. Milner, F. AbuFarha, C. Bunget, T. Kurfess, and V. H. Hammond, “Grain refinement and mechanical properties of CPTi processed by warm accumulative roll bonding,” Mater. Sci. Eng., A 561, 109–117 (2013). 3. S. Yu. Mironov, G. A. Salishchev, M. M. Myshljaev, and R. Pippan, “Evolution of misorientation distribution during warm "Abc” forging of commercialpurity tita nium,” Mater. Sci. Eng., A 418, 257–267 (2006). 4. S. V. Zherebtsov, G. S. Dyakonov, A. A. Salem, S. P. Malysheva, G. A. Salishchev, and S. L. Semiatin, “Evolution of grain and subgrain structure during cold rolling of commercialpurity titanium,” Mater. Sci. Eng., A 528, 3474–3479 (2011). 5. S. P. Malysheva, G. A. Salishchev, R. M. Galeev, V. N. Danilenko, M. M. Myshlyaev, and A. A. Popov, “Changes in the structure and mechanical properties of submicrocrystalline titanium during deformation in a temperature range of (0.15–0.45)Tm,” Phys. Met. Metallogr. 95, 390–397 (2003). 6. W. Pachla, M. Kulczyka, M. SusRyszkowska, A. Mazur, and K. J. Kurzydlowski, “Nanocrystalline titanium produced by hydrostatic extrusion,” J. Mater. Proc. Tech. 205, 173–182 (2008). 7. V. V. Stolyarov, Y. T. Zhu, I. V. Alexandrov, T. C. Lowe, and R. Z. Valiev, “Grain refinement and properties of


8.

9.

10.

11.

12.

13.

14.

15.

16.

17.

18.

19. 20.

21.

319

pure Ti processed by warm ECAP and cold rolling,” Mater. Sci. Eng., A 343, 43–50 (2003). S. V. Zherebtsov, G. S. Dyakonova, A. A. Salem, V. I. Sokolenko, G. A. Salishchev, and S. L. Semiatin, “Formation of nanostructures in commercialpurity titanium via cryorolling,” Acta Mater. 61, 1167–1178 (2013). R. Valiev, “Nanostructuring of metals by severe plastic deformation for advanced properties,” Nature Mater. 3, 511–516 (2004). R. Imayev, A. Nazarov, and R. Mulyukov, “Principles of fabrication of bulk ultrafinegrained and nanostruc tured materials by multiple isothermal forging,” Mater. Sci. Forum 638–642, 1702–1707 (2010). R. Srinivasan, D. Miracle, and S. Tamirisakandala, “Direct rolling of ascast Ti–6Al–4V modified with trace additions of boron,” Mater. Sci. Eng., A 487, 541–551 (2008). S. Roy, A. Sarkar, and S. Suwas, “On characterization of deformation microstructure in boron modified Ti⎯6Al–4V alloy,” Mater. Sci. Eng., A 528, 449–458 (2010). S. Roy and S. Suwas, “The influence of temperature and strain rate on the deformation response and micro structural evolution during hot compression of a tita nium alloy Ti–6Al–4V–0.1B,” J. Alloys Compd. 548, 110–125 (2013). R. A. Gaisin, V. M. Imaev, R. M. Imaev, and E. R. Gai sina, “Microstructure and hot deformation behavior of twophase boronmodified titanium alloy VT8,” Phys. Met. Metallogr. 114, 339–347 (2013). E. R. Gaisina, R. A. Gaisin, V. M. Imaev, and R. M. Imaev, “Microstructure and mechanical charac teristics of VT10based composite material reinforced by titanium monoboride,” Phys. Met. Metallogr. 114, 623–629 (2013). A. A. Il’in, B. A. Kolachev, and I. S. Pol’kin, Titanium Alloys. Content, Structure, Properties. A Handbook (VILS–MATI, Moscow, 2009) [in Russian]. S. Roy, V. Tungala, and S. Suwas, “Effect of hypoeutec tic boron addition on the β transus of Ti6Al4V alloy,” Metall. Mater. Trans. A 42, 2535–2541 (2011). G. A. Salishchev, R. M. Galeev, and O. R. Valiakhme tov, “Dynamical recrystallization of titanium,” Izv. Ross. Akad. Nauk, Met., No. 1, 125–129 (1994). B. A. Kolachev, Physical Metallurgy of Titanium (Met allurgiya, Moscow, 1976) [in Russian]. H. Conrad, “Effect of interstitial solutes on the strength and ductility of titanium,” Prog. Mater. Sci. 26, 123– 403 (1981). E. V. Nesterova and V. V. Rybin, “Mechanical twinning and fragmentation of commercialy pure titanium on the stage of developed plastic deformation,” Fiz. Met. Metalloved. 59, 395–406 (1985).

Translated by S. Gorin

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