TiB metal-matrix composite during high-temperature compression tests

Materials and Design 112 (2016) 17–26

Contents lists available at ScienceDirect

Materials and Design journal homepage: www.elsevier.com/locate/matdes

Deformation behavior and microstructure evolution of a Ti/TiB metal-matrix composite during high-temperature compression tests M. Ozerov, M. Klimova, A. Kolesnikov, N. Stepanov, S. Zherebtsov ⁎ Belgorod State University, Belgorod, Russia

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• Ti/TiB composite was produced by spark plasma sintering at 850 (α phase field) and 1000 °C (β phase field) • Mechanical behavior and microstructure of the composite during deformation at 500–1050 °C depended on sintering temperature • Continuous dynamic recrystallization at 700 °C and discontinuous dynamic recrystallization above 850 °C developed in Ti matrix • Stable aspect ratio of TiB whiskers reaching during deformation in the α phase field was ~10

a r t i c l e

i n f o

Article history: Received 13 July 2016 Received in revised form 13 September 2016 Accepted 14 September 2016 Available online 15 September 2016 Keywords: Metal-matrix composite Titanium alloy Spark plasma sintering Deformation Microstructure evolution Deformation behavior

a b s t r a c t A Ti/TiB metal-matrix composite (MMC) was produced by spark plasma sintering (SPS) using a Ti-10wt.%TiB2 powder mixture at temperatures of 850 or 1000 °C, corresponding to the α or β phase fields of Ti, respectively. The structures of both conditions were similar; however, a higher fraction of unreacted TiB2 and finer TiB whiskers were observed after sintering at 850 °C. The mechanical behavior and microstructural response of the MMC to uniaxial compression in the temperature range from 500 to 1050°С was studied. At low temperatures of 500– 700 °C, the MMC sintered at 1000 °C had a greater strength and lower ductility. At higher temperatures of 850– 1050 °C, the MMC sintered at 1000 °C was compressed to a 70% strain without cracking, whereas surface cracks were observed in the specimen sintered at 850 °C. Microstructure evolution of the Ti matrix was associated with the i) formation of a cell microstructure at 500 °C, ii) continuous dynamic recrystallization at 700 °C, and iii) discontinuous dynamic recrystallization at temperatures ≥850 °C. A considerable decreasing the length of the TiB whiskers, which occurred much slower in the β phase, was observed during deformation. The mechanisms governing the mechanical behavior in different temperature regions were established using activation energy analysis. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction

⁎ Corresponding author. E-mail addresses: [email protected], [email protected] (S. Zherebtsov).

http://dx.doi.org/10.1016/j.matdes.2016.09.051 0264-1275/© 2016 Elsevier Ltd. All rights reserved.

Titanium and titanium alloys are attractive for various applications, including the aerospace, automotive, chemical and biomedical industries, due to their high strength-to-density ratio, excellent corrosion resistance, and good biocompatibility. However, some applications of

18

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

titanium alloys and, particularly, commercial-purity Ti are limited because of low strength and hardness values. As one of the most bioinert metallic materials, non-magnetic commercial-purity Ti with improved strength-related properties can be used, for example, for the production of surgical cutting and clamping instruments or dental and orthopedic implants. There are some effective ways to improve the mechanical properties of Ti and Ti alloys by the inserting ceramic fibers or particles into the Ti matrix to create a metal-matrix composite (MMC), such as laser 3D printing [1,2] or spark plasma sintering (SPS) [3,4]. The former method is suitable for the surface treatment whereas the latter appears more attractive for the production of bulk specimens. Among many possible reinforcements, TiB appears to be the most attractive due to its excellent thermodynamic stability and having similar thermal expansion coefficients as the Ti matrix [3,5]. During the SPS process, TiB can be obtained in situ through the reaction of TiB2 + Ti → 2TiB, which produces a clean interface and excellent interface bonding between the TiB whiskers and the Ti matrix [3,4]. Furthermore, along with the increased strength, Ti/TiB MMCs demonstrate very low room-temperature ductility, which is inadmissible for structural materials. Hot or warm working of the MMCs can result in some increase in the low-temperature ductility [6,7,8,9]. However, Ti/TiB MMCs are more difficult to process compared with other classes of titanium alloys. Therefore, comprehensive investigations of the high-temperature deformation characteristics of Ti/TiB composites and the microstructural response to warm/hot working are still needed to determine the optimal conditions of the hot-deformation processing. It is obvious that all properties of the MMCs are dependent on their structure and, thus, on the sintering conditions. One of the most important parameters is the sintering temperature [10]. In the case of Ti-matrix composites, the polymorphism of pure Ti has to be taken into account during the selection of the sintering temperature. It is well known that the TiB whiskers have an orientation relationship (OR) with the Ti matrix [4]. However, due to the lower symmetry of the HCP α-Ti in comparison with the BCC β-Ti, the mutual phase orientation in the final product can depend on the SPS temperature. In the case of sintering at temperatures greater than the α ↔ β transition (in the β phase field) and further cooling to the α phase field, some boundaries can deviate from the OR [11,12]. Therefore, this fact can have some effect on the mechanical behavior of the composite. However, the effect of the polymorphic transformation exhibited by the Ti matrix of the Ti/TiB MMCs on the structure and mechanical behavior has not yet been fully explored. Therefore this study was focused on determining the deformation characteristics and microstructure evolution of Ti/TiB MMC produced by SPS at temperatures corresponding to either the α or the β phase fields of the titanium matrix (850 or 1000 °C, respectively) at several temperatures and strain rates.

2. Materials and procedure Commercial Ti powders (wt.% of impurities: 0.07 N, 0.05C, 0.34H, 0.34 (Fe + Ni) and 0.1Si; the α ↔ β phase transition temperature was ~910 °C) and TiB2 (wt.% of impurities: 0.04O, 0.04C, 0.02Fe) were used as the raw materials. In both cases, the particles had an irregular shape; the average particles size of the Ti and TiB2 powders were ~25 and ~ 4 μm, respectively (Fig. 1). A mixture of 90 wt.% Ti and 10 wt.% TiB2 (which yields 17 vol.% of TiB during the reaction of TiB2 + Ti → 2TiB [13]) was prepared using a Retsch RS 200 vibrating cup mill for 1 h in ethanol at the milling rotation speed of 700 rpm. The amount of TiB2 was selected based on our own preliminary results as well as literature data [3] to attain a high strength in combination with some (very low though) ductility. The room temperature ductility can supposedly be further increased using thermo-mechanical processing [6]. The Ti/TiB metal-matrix composite was produced using the SPS process under vacuum using a Thermal Technology SPS 10–3 at 850 °C (α phase field) or 1000 °C (β phase field) at 40 MPa for 15 min. Cylindrical specimens measuring 19 mm in diameter × 15 mm were obtained. Cylindrical compression samples measuring 7 mm in diameter × 10 mm were machined from the sintered Ti/TiB preforms. Specimens were compressed isothermally in air at 500, 700 or 850 °C (α phase field) and 950, 1000 or 1050°С (β phase field) in an Instron mechanical testing machine at a nominal strain rate of 10−3 s−1 to an average true strain ε of 1.2 (corresponding to a 70% reduction). The local effective strains at the center of the samples (at which all metallographic measurements and observations were made) were larger than the average sample deformation, however [14]. Strain rate jump compressive tests were conducted at temperatures 800, 850, 950, 1000 and 1050 °C and at strain rates 10−2, 5 × 10−3, 10−3, and 5 × 10−4 s−1 to determine the activation parameters. The phase composition of the composite was determined using a ARL-Xtra X-ray diffractometer with CuKα radiation. Quantitative phase analysis was performed using the Rietveld method [15] with PowderCell software. A JEOL JEM-2100FX transmission electron microscope (TEM) and a Quanta 600 field emission gun-scanning electron microscope (FEG-SEM) were used for microstructure examination. For SEM observations and X-ray analysis the specimens were mechanically polished in water with different SiC papers and a colloidal silica suspension; the final size of the Al 2O3 abrasive was 0.05 μm. Etching was carried out with Kroll's reagent (95% H2O, 3% HNO3, 2% HF). Thin-foil specimens for TEM characterization were prepared using mechanical thinning followed by electropolishing on a double jet TENUPOL-5 at 29 V and at − 35 °C using an electrolyte containing 60 ml perchloric acid, 600 ml methanol and 360 ml butanol. The Vickers microhardness was determined under a load of 10 g for 10 s. The reported hardness value is an average of at least 10 measurements.

Fig. 1. Particles of the starting powders: (a) CP-Ti and (b) TiB2.

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

19

3. Results 3.1. Initial microstructure

Fig. 2. XRD patterns of the Ti/TiB MMC sintered at 850 °C (a) and 1000 °C (b).

All microstructure observations focused on the central (most highly deformed) region of the compression specimens; all SEM and TEM micrographs were taken such that the compression direction was vertical. The linear intercept method (the lengths of intercepts along or across each particle) was used to determine the average length or diameter of the TiB whiskers. For this purpose, a total area of approximately 1150 μm2 was examined for each condition.

The result of the X-ray analysis showed (Fig. 2) that the Ti/TiB MMC sintered at 850 °C and 1000 °C consisted of (i) α Ti (hexagonal lattice, a = 0.29 nm, c = 0.47 nm), (ii) TiB2 (hexagonal lattice, a = 0.3 nm, c = 0.32 nm) and (iii) TiB (orthorhombic lattice, a = 0.61 nm, b = 0.30 nm; c = 0.45 nm). The diffraction maximums from the α-Ti were slightly broader after sintering at 1000 °C; the full width at half maximum (FWHM) of the well-defined (0002), (1011) and (1010) peaks was ~ 0.02° greater compared with the corresponding values after 850 °C sintering. The peak broadening indicated greater dislocation density and/or greater residual stresses. The intensity of the peaks corresponding to the TiB2 decreased with the sintering temperature, which indicates a decrease in the volume fraction of this phase. The fractions of the phases determined by the quantitative X-ray analysis were 74.0% Ti, 15.6% TiB and 10.4% TiB2 after sintering at 850 °C and 79.2% Ti, 18.7% TiB and 2% TiB2 after sintering at 1000 °C. The microstructure of the Ti/TiB MMC produced using SPS at 850 or 1000 °C consisted of TiB whiskers (some examples are marked in Fig. 3a, b by arrows) heterogeneously distributed within the Ti matrix. The diameter of the TiB whiskers varied in a wide interval from tens to few hundred nanometers. However, the average diameter of the TiB whiskers in the specimens obtained at 850 °C was found to be 36 ± 15 nm, which is approximately two times smaller than that obtained at 1000 °C, which was, 63 ± 35 nm. In addition, unreacted TiB2 particles

Fig. 3. SEM microstructure of the Ti/TiB MMCs sintered at 850 (a) and 1000 °C (b).

Fig. 4. Typical TEM images of the MMC microstructure (a) and the distribution of TiB in the Ti matrix with an inserted transverse section of a TiB whisker (b) in a specimen obtained using SPS at 1000 °C.

20

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

Fig. 5. Flow curves obtained during compression at 500, 700, 850, 950, 1000 and 1050 °C and a nominal strain rate of 10−3 s−1 of the Ti/TiB MMC sintered at 850 (a) and 1000 °C (b).

were observed in the microstructure of the specimens sintered at both temperatures. The fraction of TiB2 decreased with an increase in the sintering temperature from 850 to 1000 °C (Fig. 3), which is consistent with the X-ray diffraction results. The residual porosity was ~ 0.5% in both specimens. The microhardness of the specimens depended on the structural constituent. The maximum hardness of 943 ± 175 HV corresponded to the TiB2 particles. The Ti matrix with a small amount of TiB resulted in 642 ± 25 HV; an increase in the density of the TiB whiskers in the Ti matrix increased the microhardness to 762 ± 45 HV. Similar values of microhardness were obtained for both conditions of the MMC (i.e., sintered at 850 and 1000 °C). TEM examination showed that the microstructure of the Ti/TiB MMC consisted of separated TiB whiskers embedded into the Ti matrix (Fig. 4a). The microstructure of the MMCs is quite heterogeneous. In some places moderate-to-low dislocation density and residual stresses enabled distinguishing separate TiB whiskers (left part of Fig. 4a). However, the majority of the microstructure contained a very high dislocation density possibly due to a large number of TiB particles. The angle between the visible individual whiskers was approximately 60° suggesting an orientation relationship between the TiB and the Ti matrix (Fig. 4b). The TiB whiskers had an irregular hexagonal shape (insert in Fig. 4b) with sides formed by the (100), (101) and (101) planes. Many stacking faults were observed in the (100) plane of the TiB whisker. The SAED taken from both the TiB particle and the Ti matrix (not shown here) indeed indicated the fulfillment of the known OR between

To establish the effect of temperature on the mechanisms of microstructure evolution, the apparent activation energy of the deformation in this temperature interval were determined. To this end, the strain rate jump compression tests in the strain rate intervals of 5 × 10−4 to 10−2 s−1 at 800–1050 °C for both MMC conditions were performed. The obtained values of the yield strength σ0.2 and steady state flow stress σss are summarized in Table 1. Both the yield strength and the steady state flow stress values decreased with an increase in the testing temperature and a decrease in the strain rate (Table 1). No substantial difference between the mechanical characteristics of the specimens sintered at 850 °C or 1000 °C was found. The relation between the strain Table 1 Compressive mechanical properties of the Ti/TiB MMC at temperatures of 800–1050 °C and strain rates of 5 × 10−4–10−2 s−1. Temperature of SPS, °C Deformation temperature, ° C

Strain rate, s−1 σ0.2 σss

850

1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4 1 × 10−2 5 × 10−3 1 × 10−3 5 × 10−4

800

850

950

TiB and Ti [4] ((1010)α//(100)TiB and [0110]α//[011]TiB)). It should be noted that no visible differences were found between the specimens obtained by SPS at 850 or 1000 °C during the TEM examination, most likely due to the limited observation areas and heterogeneity of the produced microstructures.

1000

1050

3.2. Mechanical behavior 1000

The stress-strain curves for both conditions of the MMC (produced by SPS at 850 or 1000 °C) compressed in the temperature range from 700 to 1050 °C at a nominal strain rate of 10−3 s−1 exhibited an initial hardening transient, a peak flow stress, and then flow softening (Fig. 5). Such behavior can be ascribed to the dynamic restoration processes (recrystallization or recovery) [16]. During deformation at a lower temperature of 500 °C the mechanical behavior of the two studied conditions differed considerably. The MMC produced by SPS at 850 °C demonstrated continuous strengthening after the beginning of plastic flow (Fig. 5a). In contrast, the MMC produced by SPS at 1000 °C had an extended plateau (Fig. 5b). Note that at 500 °C the yield strength of the composite sintered at 1000 °C was much greater than that of the composite sintered at 850 °C. The difference became less (still being sufficient, however) at 700 °C. At higher testing temperatures, the strength of the composites was not significantly affected by the sintering conditions.

800

850

950

1000

1050

116 100 86 71 115 98 71 70 99 87 63 44 94 57 50 48 80 68 42 37 128 107 80 65 122 101 76 63 92 76 44 36 82 71 35 34 73 62 33 32

118 111 90 85 116 105 85 81 100 94 76 72 96 85 67 63 81 76 60 55 130 119 98 85 125 109 90 77 95 86 71 62 85 78 59 52 74 68 52 46

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

21

Fig. 6. Logarithmic plot of the dependence of the strain rate on the steady-state flow stress for the Ti/TiB MMC produced using SPS at 850 (a) or 1000 °C (b).

Fig. 7. Arrhenius plot of the logarithmic steady-state flow stress vs. the inverse of temperature in the interval 800–1050 °C for the Ti/TiB MMC produced using SPS at 850 (a) or 1000 °C (b).

rate and the temperature during high-temperature deformation can generally be described with the well-known Zener–Hollomon parame_ [17,18,19]. ter Z(ε) n

Zðε_ Þ ¼ Aε_ ¼ ε_ expðQ =RTÞ

deformation with a weaker temperature dependence [20]. In accordance with Eq. (3) the value of the apparent activation energy for the MMC produced by SPS at 850 °C was found to be 250 kJ/mol and

ð1Þ

where ε_ is the deformation strain rate, n is the stress exponent, Q is the activation energy, A is a constant sensitive to the deformation mechanism, R is the gas constant and T is the absolute temperature. The parameters n and Q can be determined as n¼

∂ ln ε_ ∂ ln ðσ ss Þ T

Q ¼ −R

∂ ln ε_ ∂ ln ðσ ss Þ ≡ nR ; ∂ð1=TÞ σ ∂ð1=TÞ ε_

ð2Þ

ð3Þ

where σss denotes the steady state flow stress. To determine the stress exponent n, logarithmic plots of the strain rate against the steady state stress (corresponding values were taken from Table 1) at different temperatures for both MMC conditions were plotted (Fig. 6). The experimental data for temperatures 800–1050 °C can be approximated rather well using a linear dependence with the slope n ranged from ~ 7.5 to 8.7 (Fig. 6a) and from 6.2 to 7.2 (Fig. 6b) for the MMC obtained using SPS at 850 and 1000 °C, respectively. The dependences of the logarithm of the steady state stress values on the inverse absolute temperature, 1/T, at four different strain rates for both MMC conditions are given in Fig. 7. The observed relationships can be approximated as a linear function with different slopes at temperatures in the intervals 950–1050 °C or 800–950 °C, which corresponded to the β or α phase fields, respectively. It worth noting that the slopes were almost independent of the strain rates. At a lower temperature, the slopes decreased suggesting a transition to

Fig. 8. Specimens of the Ti/TiB MMC produced using SPS at 850 °C (a, c, e, g, i, k) or 1000 °C (b, d, f, h, j, i) after isothermal compression at 500 (a, b), 700 (c, d), 850 (e, f), 950 (g, h), 1000 (i, j) and 1050 °C (k, l).

22

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

107 kJ/mol for the high-temperature and low-temperature intervals, respectively (Fig. 7a). For the MMC produced by SPS at 1000 °C, the apparent activation energy was 239 kJ/mol for the high-temperature interval and 168 kJ/mol for the lower temperatures (Fig. 7b). 3.3. Microstructure evolution during deformation The samples of Ti/TiB MMC can be successfully worked at elevated temperatures as shown previously (Fig. 5). However, the deformability of the composite depended on the temperatures of both sintering and deformation. During deformation at 500 °C, the specimens produced both at 850 and 1000 °C fractured before attaining 70% strain (Fig. 8a, b). Note that the deformability of the sample sintered at 1000 °C (Fig. 8b) was much lower than that of the sample sintered at 850 °C (Fig. 8a). Moreover, it appears that plastic flow in the specimen sintered at 1000 °C during compression at 500 °C was highly localized, which can be the reason for the steady-state flow stage observed in the stress-

strain curve (Fig. 5b). At 700 °C, the specimens were compressed to 70%, but deep cracks formed at the side faces (Fig. 8c, d). The formation of the surface cracks was observed during deformation at higher temperatures (till 1050 °C) of the specimens sintered at 850 °C (Fig. 8e, g, i, k). Furthermore, the samples sintered at 1000 °C and then compressed to 70% in the interval 850–1050 °C did not have any visible surface cracks (Fig. 8f, h, j, l). Typical microstructures of the specimens compressed to 70% at temperatures 500, 700, 850, 950, 1000 and 1050°С and a nominal strain rate of 10−3 s−1 are shown in Fig. 9 with respect to the MMC obtained by SPS at 850 °C. The most prominent effect of the decreased deformation temperature on the microstructure was associated with more intensive rearrangement of the TiB whiskers along the metal flow direction. As a result of the deformation at lower temperatures (500–950 °C), welldefined bands formed by the TiB whiskers were formed (Fig. 9a–d). At higher temperatures the TiB whiskers remained quite heterogeneously distributed within the Ti matrix after deformation. In addition, the

Fig. 9. SEM microstructure of the Ti/TiB MMC produced using SPS at 850 °C; uniaxial compression at a nominal strain rate of 10−3 s−1 and temperatures 500°С (a), 700°С (b), 850°С (c), 950°С (d), 1000°С (e) and 1050°С (f). The compression axis is vertical in all cases.

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

fraction of the TiB2 remnant decreased with an increase in the deformation temperature. The microstructure of the Ti matrix cannot be recognized in these SEM images. Specimens obtained by SPS at 1000 °C demonstrated a rather similar dependence of the microstructure on the deformation temperature. The TEM analysis showed the obvious dependence of the microstructure on the deformation temperature. The microstructure of the MMCs after compression to 70% at 500 °C 10−3 s−1 was found to be typical of the largely deformed condition of titanium [21] and consisted of a cellular structure with a high dislocation density (Fig. 10a). The boundaries of the cells were rather wide and loose. The size of the cells varied over a wide range from a hundred to a few hundred nanometers. The interphase Ti/TiB boundaries are blurred due to high internal stresses caused by the high dislocation density in the vicinity of the interfaces. Deformation at a temperature of 700 °C resulted in the formation of areas that were 1–1.5 μm in diameter with a much lower dislocation density compared with the neighboring fields (Fig. 10b). The boundaries of those dislocation-free areas were sinuous and unclear. However at a higher temperature of 850 °C, the formation of a partially recrystallized microstructure was observed (Fig. 10c). The size of the recrystallized grains with well-defined grain boundaries was found to be ~1 μm. More pronounced dynamic recrystallization with the formation of a perfect grain structure occurred at 1000 °C; the size of the grains was found to be ~2.5 μm (Fig. 10d). The development of the dynamic recrystallization was more evident in the areas of the Ti matrix with a relatively low density of the TiB whiskers. The greater density of the TiB particles hindered grain boundary movement thereby preserving deformed microstructure [16]. After deformation at temperatures from 700 to 1000 °C, the Ti/TiB interfaces were clear without any sign of defects (pores or cracks), which could indicate deformation incompatibility. It is worth noting that the TEM examination did not reveal any noticeable difference in

23

the microstructure of the deformed specimens, which had been obtained by SPS either at 850 or at 1000 °C. Quantitative analysis showed a somewhat similar effect of the deformation temperature on the diameter and the apparent length of the TiB whiskers for specimens obtained using SPS at both 850 and 1000 °C (Fig. 11). The diameter (Fig. 11a) and the apparent length (Fig. 11b) of the whiskers did not almost change during deformation at lower temperatures (500–850 °C) for both (SPS at 850 °C and SPS at 1000 °C) conditions. At higher temperatures, the diameter and the apparent length continuously increased for both conditions achieving values approximately 1.5 or 3 times greater than that at 500 °C for SPS at 1000 or 850 °C, respectively. The faster kinetics of the TiB whisker coarsening at temperatures N850 °C was most probably due to the decomposition of the retained TiB2 and the increased diffusivity of boron to the Ti/TiB interfaces. The absolute values of the apparent length of the TiB whiskers were always smaller in the case of SPS at 850 °C in comparison with those for SPS at 1000 °C: ~ 2 times greater after deformation at temperatures b850 °C and ~1.2 greater at higher temperatures. 4. Discussion The results obtained from the current study suggest that the mechanical behavior, the microstructure evolution during compression in the temperature interval 500–1050 °C and the quality of the Ti/TiB MMCs depends significantly on the SPS temperature. Obviously, the reasons for the differences should be related to the parameters of the initial (as-sintered) structure of the composites. One of the parameters that can significantly influence the deformability of the MMCs is the volume fraction of the residual TiB2 the content of which decreases when increasing the temperature of the SPS. The increase in the SPS temperature from 850 °C (α phase field) to 1000 °C (β phase field) decreased the XRD-determined fraction of TiB2 from 10.4 to 2%. The excessive amount

Fig. 10. Typical TEM images of the Ti/TiB MMC obtained using SPS at 1000 °C after uniaxial compression to 70% at 500 (a), 700 (b), 850 (c) or 1000 °C (d) and a strain rate of 10−3 s−1.

24

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

Fig. 11. Diameter (a) or apparent length (b) of the TiB whiskers in Ti/TiB MMCs produced using SPS at 850 or 1000 °C after 70% strain as a function of the deformation temperature.

of the very hard TiB2 can result in pronounced cracking of specimens that were spark plasma sintered at 850 °C even at elevated deformation temperatures (Fig. 8). Some decrease in the ductility and an increase in strength can be caused by finer TiB reinforcements in the case of the MMC produced using SPS at 850 °C (~30 nm vs. ~ 60 nm in the MMC produced by SPS at 1000 °C). Furthermore, it was reported in [22] that the aspect ratio of the TiB (rather than the diameter) is the main factor that influences the mechanical properties of the TMCs. In the initial state, the aspect ratio of the TiB whiskers was ~ 47 and ~ 55 for MMCs produced using SPS at 1000 or 850 °C, respectively, which should give rather similar strengths. However, Fig. 5 shows that at 500 and 700 °C the MMC produced at 1000 °C is even stronger thereby suggesting the presence of some other reasons for the different strengths and ductility of the MMCs produced at different temperatures. One more factor that can influence the macroscopic properties of the specimens obtained at different SPS temperatures is variations in the structure of the Ti/TiB interfaces due to the presence of the α↔β polymorphic transformation. When TiB whiskers form in the β phase field (SPS conducted above ~910 °C), the crystal lattices of the TiB and β phases satisfy the (110)β//(001)TiB and [111]β//[010]TiB OR. Cooling down to temperatures b 910 °C results in the β → α transition, which thereby decreases the possible variant selections because of the lower symmetry of the HCP α in comparison with the BCC β. The classical Burgers OR ((0001)α//(110)β and [1120]α//[111]β) [23] gives 6 variant selections when the α phase forms in the β phase at the β → α transition. The {110} family plane in the β phase is the habit plane for both the α phase and the TiB particles. Therefore, there is an obvious

α/TiB mismatch are likely relaxed prior to deformation, and therefore, little difference in the mechanical behavior is observed. The results of the activation energy analysis (Fig. 7) suggest that the two temperature intervals with different Q values exist at temperatures ≥850 °C. The observed change of Q with the deformation temperature is most likely associated with the α↔β phase transition of the Ti matrix at ~910 °C. The obtained values of the apparent activation energy in the top of the α phase field (107–168 kJ/mol) are consistent with the reported activation energies of bulk diffusion in the α phase (150 kJ/mol) of titanium [24,25]. In the case of the β phase field, the values of the apparent activation energy (Q = 239–250 kJ/mol) are greater than those observed earlier for self-diffusion in titanium (131 to 153 kJ/mol) [24,25], which is, however, in better agreement with the values of Q reported for the Ti/TiB MMCs (209 kJ/mol) [26]. Greater values of the activation energy and stress exponent in the case of the composite can be ascribed to effective inhibition of the movement of dislocations by the TiB reinforcements [26]. Some increase in the Q values compared with the literature data can also be associated with the development of dynamic recrystallization during deformation [27]. The value of strain rate sensitivity for both conditions m = 1/n ≈ 0.12 (Fig. 10) is consistent with the results shown in [28] and, together with the values of Q, implies domination of dislocation-related mechanisms (dislocation glide/climb) in the interval 800–1050 °C [29]. Furthermore, the evaluation of the activation energy at lower temperatures (i.e., in the interval 500–850 °C) and a strain rate of 10−3 s− 1 using n = 7 gave a noticeably greater activation energy (Q ≈ 300 kJ/mol, Fig. 12) than that obtained in the top of the α phase

OR between the α phase and the TiB: (0001)α//(001)TiB and [1120]α// [010]TiB (more usual form of the OR that can be found in literature [4] is (1010)α//(100)TiB and [0110]α//[011]TiB). However, if the habit planes for both the α phase and TiB are not exactly the same (still belonging the {110}β family) a strict OR between the phases can be not held. For example, if the habit planes are (110)β for the ТiB and (011)β for the α phase, there is no corresponding parallel low-index planes in the α and TiB. The closest to parallelism planes were found to be (1101)α and (001)ТiB with a 1.4° angle between them. The mismatches between the TiB and α Ti were most likely accommodated by the more ductile Ti matrix, which should result in the increased Ti lattice distortion and possibly a greater dislocation density. This suggestion is in agreement with the previously noted broadening of the Ti peaks in the XRD patterns obtained from the MMC sintered at 1000 °C. In the case of the specimen that did not undergo the β → α transition (i.e., which was produced at 850 °C), no deviation from the OR between α/TiB is expected. Therefore it is hypothesized that the possible deviations from the Burgers OR in the specimens sintered at 1000 °C are responsible for their greater strength at 500 °C or 700 °C and the lower ductility at 500 °C. At temperatures of ≥ 850 °C, the distortions generated by the

Fig. 12. Arrhenius plot of the logarithmic steady-state flow stress vs. the inverse of the temperature for a temperature interval 500–1050 °C and a strain rate of 10−3 s−1.

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

field (Fig. 11). This result is in good accordance with the values of Q = 250–330 kJ/mol reported in [29] for glide on the first order prism planes when the principal rate controlling mechanism is the thermally activated overcoming of the interstitial solute atom obstacles. The interstitial solutes in the present case were likely boron atoms dissolved into the Ti matrix during in situ sintering. The analysis of the microstructures of the deformed MMCs indicates a change in the mechanism of microstructure formation in the low-temperature interval that also can result in some change of the activation energy. Continuous strengthening and cellular microstructure with very high dislocation density, which is typical of largely deformed condition, were observed after deformation at 500 °C and was replaced by partially recrystallized structure at 700 °C (Fig. 10a, b). The development of recrystallization at this temperature is consistent with a characteristic for the dynamic recrystallization mechanical behavior of the MMC (Fig. 5) when the stress-stain curves show a pronounced steady-state flow stage. At low temperatures, the formation of the new grains occurs due to an increase in the subgrain misorientations as a result of the progressive accumulation of dislocations in subboundaries, i.e., continuous dynamic recrystallization (cDRX) [16, 19]. Large areas of the Ti matrix with low dislocation densities surrounded by diffuse dislocation boundaries can suggest the occurrence of cDRX. In contrast to low temperatures, the microstructure evolution during deformation at 850 °C and greater appears to be controlled by discontinuous dynamic recrystallization (dDRX) [16,19] that is associated with the conventional nucleation and growth of new recrystallized grains (Fig. 10c, d). The change in the mechanisms of the microstructure evolution from cDRX to dDRX with an increase in the deformation temperature is well established for many metallic systems including commercially pure titanium [30]; however, in the case of MMSc, this question deserves further investigations. The deformation characteristics of the Ti matrix controlled the response of the TiB particles on the plastic working. Deformation in the studied temperature interval was accompanied by a considerable decrease in the length of the TiB whiskers. However, in the β phase field, this shortening occurred much slower (especially at higher temperature) compared with the α phase field (Fig. 11b). Due to the greater softness and ductility of the β phase compared with the α phase, hard TiB whiskers can maintain a high aspect ratio during deformation at temperatures N910 °C. In the α phase field, the involvement of the TiB particles into deformation increases and the apparent length of TiB whiskers achieved the minimum values. The diameter of the whisker, in contrast, is weakly affected by deformation in the α phase field and increased significantly with an increase in the temperature above the α↔ β phase transition (Fig. 11a). A pronounced increase in the thickness of the TiB at temperatures N 850–950 °C is likely associated with a high rate of diffusion in the β phase field. It is worth noting that the apparent length and diameter of the TiB whiskers did not change when decreasing the temperature in the α phase field, which suggests achieving a stable aspect ratio of ~10 for both conditions of the MMCs (produced by SPS at 850 or 1000 °C). 5. Conclusions The mechanical behavior and microstructural response of a Ti/TiB metal-matrix composite produced using spark plasma sintering (SPS) with a Ti-10wt.%TiB2 powder mixture at 850 or 1000 °C to uniaxial compression in the temperature interval from 500 to 1050°С was studied. The following conclusions can be drawn from this work: 1. The deformability and mechanical properties of the Ti/TiB MMC depend on the temperature at which the MMC was produced using SPS. The MMC sintered at 1000 °C (higher than the α ↔ β phase transition) showed greater strength at low temperatures (500 and 700 °C) and deformed without cracks at temperatures ≥850 °C in contrast to

25

the MMC sintered in the top of the α phase field. The greater strength of the MMC produced by SPS at 1000 °C in comparison to the MMC produced using SPS at 850 °C can be ascribed to a greater residual stress that arises from the deviation from the orientation relationship due to the α ↔ β phase field transition during cooling from the sintering temperature. 2. Microstructure evolution of the Ti matrix for both MMC conditions (i.e., produced using SPS at 850 or 1000 °C) is associated with the i) formation of a cell microstructure with high dislocation density at 500 °C, ii) continuous dynamic recrystallization at 700 °C, and iii) discontinuous dynamic recrystallization at temperatures N850 °C. The change in the microstructure evolution mechanism was associated with a decrease in the apparent activation energy from Q ≈ 300 kJ/mol for the interval 500–800 °C to Q = 107–168 kJ/mol for T = 800–950 °C. Another change in the apparent activation energy to Q = 239–250 kJ/mol was caused by the α↔β phase field transition. The values of the stress exponent for both conditions (n ≈ 7 for SPS 1000 °C and n ≈ 8 for SPS 1000 °C) and the apparent activation energy implied domination of the dislocation-related mechanisms (dislocation glide/climb) controlled by pipe diffusion during deformation in the interval 800–1050 °C. The increase in the apparent activation energy to ~300 kJ/mol in the low-temperature domain can be ascribed to glide on the first order prism planes when the principal rate controlling mechanism is the thermally activated overcoming of interstitial solute atom obstacles. 3. Deformation in the studied temperature interval was accompanied by the rearrangement of the TiB whiskers along the metal flow direction and a considerable decrease in the length of the TiB whiskers; both processes were found to be less pronounced in the β phase field due to greater softness and ductility of the BCC β in comparison with the HCP α. A stable aspect ratio of the TiB whiskers which was reached during deformation in the α phase field was found to be ~10 for the both MMCs conditions. Acknowledgments The authors gratefully acknowledge the financial support from the Russian Science Foundation [Grant number 15-19-00165]. The authors are grateful to the personnel of the Joint Research Centre, Belgorod State University for their assistance with the instrumental analysis. References [1] C.Y. Yap, C.K. Chua, Z.L. Dong, Z.H. Liu, D.Q. Zhang, L.E. Loh, S.L. Sing, Review of selective laser melting: Materials and applications, Appl. Phys. Rev. 2 (2015), #041101 http://dx.doi.org/10.1063/1.4935926. [2] X. Wu, J. Liang, J. Mei, C. Mitchell, P.S. Goodwin, W. Voice, Microstructures of laserdeposited Ti-6Al-4V, Mater. Design 25 (2004) 137–144, http://dx.doi.org/10.1016/j. matdes.2003.09.009. [3] K. Morsi, V.V. Patel, Processing and properties of titanium–titanium boride (TiBw) matrix composites — a review, J. Mater. Sci. 42 (2007) 2037–2047, http://dx.doi. org/10.1007/s10853-006-0776-2. [4] H. Feng, Y. Zhou, D. Jia, Q. Meng, J. Rao, Growth mechanism of in situ TiB whiskers in spark plasma sintered TiB/Ti metal matrix composites, Cryst. Growth Des. 6 (2006) 1626–1630, http://dx.doi.org/10.1021/cg050443k. [5] K.S. Ravi Chandran, K.B. Panda, S.S. Sahay, TiBw-reinforced Ti composites: processing, properties, application, prospects, and research needs, JOM 56 (2004) 42–48, http://dx.doi.org/10.1007/s11837-004-0127-1. [6] V. Imayev, R. Gaisin, E. Gaisina, R. Imayev, H.-J. Fecht, F. Pyczak, Effect of hot forging on microstructure and tensile properties of Ti-TiB based composites produced by casting, Mater. Sci. Eng. A609 (2014) 34–41, http://dx.doi.org/10.1016/j.msea. 2014.04.091. [7] H.T. Tsang, C.G. Chao, C.Y. Ma, Effects of volume fraction of reinforcement on tensile: and creep properties of in-situ TiB/Ti MMC, Scr. Mater. 37 (1997) 1359–1365, http://dx.doi.org/10.1016/S1359-6462(97)00251-0. [8] O.M. Ivasishin, R.V. Teliovych, V.G. Ivanchenko, S. Tamirisakandala, D.B. Miracle, Processing, microstructure, texture, and tensile properties of the Ti-6Al-4V-1.55B eutectic alloy, Metall. Mater. Trans. A 39 (2008) 402–416, http://dx.doi.org/10.1007/ s11661-007-9425-x. [9] C. Zhang, F. Kong, S. Xiao, H. Niu, L. Xu, Y. Chen, Evolution of microstructural characteristic and tensile properties during preparation of TiB/Ti composite sheet, Mater. Design. 36 (2012) 505–510, http://dx.doi.org/10.1016/j.matdes.2011.11.060.

26

M. Ozerov et al. / Materials and Design 112 (2016) 17–26

[10] X. Shen, Z. Zhang, S. Wei, F. Wang, S. Lee, Microstructures and mechanical properties of the in situ TiB-Ti metal–matrix composites synthesized by spark plasma sintering process, J. Alloys Compd. 509 (2011) 7692–7696, http://dx.doi.org/10.1016/j. jallcom.2011.05.017. [11] D.X. Li, D.H. Ping, Y.X. Lu, H.Q. Ye, Characterization of the microstructure in TiBwhisker reinforced Ti alloy matrix composite, Mater. Lett. 16 (1993) 322–326, http://dx.doi.org/10.1016/0167-577X(93)90201-8. [12] A. Genc, R. Banerjee, D. Hill, H.L. Fraser, Structure of TiB precipitates in laser deposited in situ Ti-6Al-4V-TiB composites, Mater. Lett. 60 (2006) 859–863, http://dx.doi. org/10.1016/j.matlet.2005.10.033. [13] L.J. Huang, L. Geng, B. Wang, L.Z. Wu, Effects of volume fraction on the microstructure and tensile properties of in situ TiBw/Ti6Al4V composites with novel network microstructure, Mater. Design. 45 (2013) 532–538, http://dx.doi.org/10.1016/j. matdes.2012.09.043. [14] N. Stefansson, S.L. Semiatin, D. Eylon, The kinetics of static globularization of Ti-6Al4V, Metall. Mater. Trans. A 33 (2002) 3527–3534, http://dx.doi.org/10.1007/ s11661-002-0340-x. [15] G. Will, Powder Diffraction: The Rietveld Method and the two-Stage Method to Determine and Refine Crystal Structures from Powder Diffraction Data, Springer, Berlin, 2005. [16] F. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Elsevier, Oxford, 2004. [17] C.M. Sellars, Recrystallization of metals during hot deformation, Phil. Trans. R. Soc. London A 288 (1978) 147–158http://www.jstor.org/stable/74980. [18] T. Sakai, J.J. Jonas, Dynamic recrystallization: mechanical and microstructural considerations, Acta Metall. 32 (1984) 189–209, http://dx.doi.org/10.1016/00016160(84)90049-X. [19] T. Sakai, A. Belyakov, R. Kaibyshev, H. Miura, J.J. Jonas, Dynamic and post-dynamic recrystallization under hot, cold and severe plastic deformation conditions, Prog. Mater. Sci. 60 (2014) 130–207, http://dx.doi.org/10.1016/j.pmatsci.2013.09.002. [20] A. Belyakov, H. Miura, T. Sakai, Dynamic recrystallization under warm deformation of a 304 type austenitic stainless steel, Mater. Sci. Eng. A 255 (1998) 139–147, http://dx.doi.org/10.1016/S0921-5093(98)00784-9.

[21] S.V. Zherebtsov, G.S. Dyakonov, A.A. Salem, V.I. Sokolenko, G.A. Salishchev, S.L. Semiatin, Formation of nanostructures in commercial-purity titanium via cryorolling, Acta Mater. 61 (2013) 1167–1178, http://dx.doi.org/10.1016/j.actamat. 2012.10.026. [22] M.Y. Koo, J.S. Park, M.K. Park, K.T. Kim, S.H. Hong, Effect of aspect ratios of in situ formed TiB whiskers on the mechanical properties of TiBw/Ti-6Al-4V composites, Scr. Mater. 66 (2012) 487–490, http://dx.doi.org/10.1016/j.scriptamat.2011.12.024. [23] W.G. Burgers, On the process of transition of the cubic-body-centered modification into the hexagonal-close-packed modification of zirconium, Physica 1 (1934) 561–586, http://dx.doi.org/10.1016/S0031-8914(34)80244-3. [24] H.J. Frost, M.F. Ashby, Deformation-Mechanism Maps: The Plasticity and Creep of Metals and Ceramics, Pergamon Press, Oxford, 1982. [25] N.E. Walsöe De Reca, C.M. Libanati, Autodifusion de titanio beta y hafnio beta, Acta Metall. 16 (1968) 1297–1305, http://dx.doi.org/10.1016/0001-6160(68)90010-2. [26] Y. Zhang, L. Huang, B. Liu, L. Geng, Hot deformation behavior of in-situ TiBw/Ti6Al4V composite with novel network reinforcement distribution, Trans. Nonferrous Metals Soc. China 22 (2012) s465–s471, http://dx.doi.org/10.1016/S1003-6326(12)617476. [27] S.V. Raj, T.G. Langdon, Creep behavior of copper at intermediate temperatures – I. Mechanical characteristics, Acta Metall. 37 (1989) 843–852, http://dx.doi.org/10. 1016/0001-6160(89)90011-4. [28] S. Kumari, N.E. Prasad, K.S.R. Chandran, G. Malakondaiah, High-temperature deformation behavior of Ti-TiBw in-situ metal-matrix composites, JOM 56 (2004) 51–55, http://dx.doi.org/10.1007/s11837-004-0129-z. [29] H. Conrad, Effect of interstitial solutes on the strength and ductility of titanium, Prog. Mater. Sci. 26 (1981) 123–403, http://dx.doi.org/10.1016/0079-6425(81)90001-3. [30] A. Belyakov, S. Zherebtsov, G. Salishchev, Three-stage relationship between flow stress and dynamic grain size in titanium in a wide temperature interval, Mater. Sci. Eng. A 628 (2015) 104–109, http://dx.doi.org/10.1016/j.msea.2015.01.036.