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Ambient-temperature nanoindentation creep in ultrafine-grained titanium processed by ECAP
Materials Science & Engineering A 676 (2016) 73–79
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Ambient-temperature nanoindentation creep in ultrafine-grained titanium processed by ECAP Xiaoyan Liu n, Qianqian Zhang, Xicheng Zhao, Xirong Yang, Lei Luo School of Metallurgical Engineering, Xi’an University of Architecture and Technology, Xi’an 710055, China
art ic l e i nf o
a b s t r a c t
Article history: Received 12 April 2016 Received in revised form 21 August 2016 Accepted 27 August 2016 Available online 29 August 2016
Ultrafine-grained (UFG) Ti was prepared by equal channel angular pressing (ECAP) through 1 and 4 passes at room temperature. The ambient-temperature creep behavior and mechanism of UFG Ti were studied by nanoindentation creep tests at the loading strain rate of 0.005 and 0.1 s 1. The effect of grain size on creep behavior and relevant creep parameters such as steady creep strain rate and creep stress exponent (n) was estimated for coarse-grained (CG) and UFG Ti. The results show that the creep resistance of UFG Ti is enhanced with respect to CG Ti. UFG Ti after 1 pass of ECAP exhibits the highest creep resistance. The creep resistance of UFG Ti decreases with increasing the number of ECAP passes. The creep stress exponents of UFG Ti are much higher than those of CG Ti while steady creep strain rates do not vary much with grain size. The creep stress exponents of CG and UFG Ti are dependent on the loading strain rate and increase with increasing the loading strain rate. The power-law creep with a stress exponent of 18.1–24.6 is consistent with dislocation process, especially for dislocation emission and annihilation at grain boundaries, which is a key mechanism in creep deformation of UFG Ti. Grain boundaries may play an important role in the creep behavior of UFG Ti. & 2016 Elsevier B.V. All rights reserved.
Keywords: Ultrafine-grained titanium Nanoindentation Creep Dislocations
1. Introduction Titanium and its alloys with high strength and excellent corrosion resistance are widely applied in the fields of aerospace and transport. Ultrafine-grained Ti processed by equal channel angular pressing (ECAP) has been under intensive attention worldwide for past two decades due to its low weight, excellent mechanical properties, and good biocompatibility for potential engineering applications and biomedical implants. To date, a large amount of studies have been reported on the microstructure evolution, deformation behavior and fatigue failure mechanism of UFG Ti [1–3]. Considering perspectives of engineering applications, the creep properties are often of primary importance. Ambient-creep was observed in coarse-grained (CG) titanium about the 1950s [4,5]. Since that time, many studies of ambienttemperature creep of titanium and its alloys have been carried out by materials researchers [6–8]. It was demonstrated that ambienttemperature creep mechanisms of Ti alloys were interpreted as the straightly aligned dislocations slip. Later, Neeraj et al. [9] and Hasija et al. [10] found that ambient-temperature creep behavior of Ti–6Al complied with the Andrade's creep law by experiments and computer simulation, respectively. n
Corresponding author. E-mail address: [email protected] (X. Liu).
http://dx.doi.org/10.1016/j.msea.2016.08.111 0921-5093/& 2016 Elsevier B.V. All rights reserved.
The grain size is expected to play an important role in creep behavior of UFG materials. However, the study on the creep behavior of UFG materials processed by ECAP was scarcely carried out. These reports concerning the creep behavior of ECAP materials were focused on pure aluminum and pure copper [11–14]. It was found that the creep resistance of the UFG materials processed by ECAP was obviously improved with respect to CG materials and the creep behavior of UFG materials strongly depends on the number of ECAP passes [14]. Only very limited results are available describing the creep behavior of UFG Ti processed by ECAP at elevated temperature [15]. The results showed that the high content of grain boundaries and dislocations in the initial state of UFG Ti enhanced its deformation resistance and stabilized the materials. The occurrence of dynamic strain aging in UFG Ti confirmed that the creep mechanism of UFG Ti was related to dislocation motion, such as dislocation slip and climb. Recently, there has been intensive interest to investigate creep behavior and mechanism of micro/nano structural materials and thin films by nanoindentation creep test [16–18]. Although the creep behavior is considerably more complex in nanoindentation creep tests than that in conventional uniaxial tensile/compression tests, it has been found that the creep stress exponent (n) can be calculated from the holding stage in nanoindentation test and the values of stress exponent are generally in good agreement with those obtained by conventional uniaxial tensile/compression tests [19,20]. The purpose of this study is to reveal the ambient-
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temperature creep behavior and mechanism in UFG Ti processed by ECAP using nanoindentation tests.
2. Experimental 2.1. Materials Ultrafine-grained Ti used in this work was prepared by ECAP following route C for 1 and 4 passes at room temperature, respectively. The more details of the processing method were reported elsewhere [21]. Its chemical composition (in wt%) is 0.012 Fe, 0.022 C, 0.003 N, 0.001 H, 0.06 O and balance Ti. The reference coarse-grained (CG) Ti was prepared by annealing the hot rolled pure Ti at 600 °C for 40 min and then air cooling. Microstructure and grain sizes of three materials were examined by transmission electron microscopy (TEM) and optical microscopy (OM). 2.2. Nanoindentation test In order to acquire reliable nanoindentation data, the surface of the indentation specimens was metallographically ground with a
series of progressively finer SiC sandpapers, and mechanically polished to mirror finishing. Nanoindentation tests were performed using the Nano indenter’ XP system (MTS, Inc.) at room temperature. The continuous stiffness measurement (CSM) system was used in this study, which provides a better control on the contact step of the indentation. Upon calibration on standard fused silicon, the tip of the Berkovich diamond indenter was estimated to have a radius of 50 nm. Loads (P) and Displacements (h) were measured with a resolution of 50 nN and 0.01 nm, respectively. All the indents were carried out to a depth limit of 1000 nm at the loading strain rate of 0.005 s 1 and 0.1 s 1, respectively. The indentation direction was parallel to the ECAP axis. In the loading stage, hardness was automatically recorded as a continuous function of surface penetration depth by the CSM mode. After the indenter reached the depth limit, the corresponding load was held for 300 s to monitor the variation in creep displacement with time. Then, the indenter was unloaded to 10% of the maximum load and held for 100 s to correct thermal drift. Finally, the indenter was withdrawn from the specimen surface to terminate the creep test. At least eight indentation tests were carried out at each loading strain rate. Indent surface morphologies were imaged in a ZEISS Supra55 FEG SEM at 15 kV using secondary electron (SE) signals.
Fig. 1. Microstructure of (a) CG Ti, (b) UFG Ti (1P), and (c) UFG Ti (4P).
X. Liu et al. / Materials Science & Engineering A 676 (2016) 73–79
3. Results and discussions 3.1. The microstructures of CG and UFG Ti The microstructures of CG and UFG Ti are shown in Fig. 1. The OM image shows that CG Ti exhibits mostly equiaxed grains with an average grain size of about 26 mm. The micrographs in Figs. 1b and c were taken from the Y plane of ECAP billets. Based on the orthogonal notation used in many earlier studies of ECAP [22], the Y plane is parallel to the side face of the billet at the point of exit from the die. For the UFG Ti after 1 pass of ECAP, elongated parallel shear bands with average widths of about 300 nm are observed in Fig. 1b. Within these parallel bands there are some areas having a high density of dislocations. The TEM image in Fig. 1c shows that the microstructure of UFG Ti after 4 passes of ECAP consists of equiaxed ultrafine grains with high-angle grain boundaries having an average grain size of about 170 nm. Selected area diffraction (SAD) pattern inserted into Fig. 1c also provides evidence for the extent of microstructure refinement. 3.2. Load-displacement curves Fig. 2 shows the representative load-displacement (P-h) curves for CG and UFG Ti at different loading strain rates. In the loading stage, the load required to reach the given displacement increases with increasing the loading strain rate for the CG and UFG Ti. This indicates a pronounced strain rate sensitivity of hardness or stress of three materials. This observation is in agreement with other findings on UFG Ti and Cu [23,24]. At the same loading strain rate, the holding loads of three materials are the highest and lowest for UFG Ti after 4 passes of ECAP and CG Ti, respectively. Furthermore, the hardness values are found to increase with decreasing the grain size, which due to the significant grain refinement and high
75
density of dislocations in UFG Ti after ECAP. Significant load plateaus are observed in P-h curves for CG and UFG Ti. Such load plateau demonstrates a significant creep deformation in the holding stage. The width of the load plateaus increases with increasing the loading strain rate for CG and UFG Ti. 3.3. Indentation morphology Representative SEM indentation micrographs of CG and UFG Ti at the loading strain rate of 0.1 s 1 are shown in Fig. 3. It can be seen that no cracking in any of the indentations, which suggests fully plastic flow in CG and UFG Ti. The indentation morphology of CG Ti (Fig. 3a) shows a fairly uniform surface topography with no visible pile-up. In contrast, the pile-up behavior of UFG Ti is observed in indentation morphology (Fig. 3b and c), indicating inhomogeneous or localized plastic flow. The dimensional analysis and finite element calculations showed that the pile-up and sinkin phenomena were determined by the value of Y/E (where E is the Young's modulus and Y is the initial yield stress), as well as the work-hardening exponent, n [25,26]. Pile-up is significant only when the value of Y/E is very small (Y/E o0.01) and the materials with a small work-hardening exponent (n o0.1) [26]. It is well known that very little work-hardening occurs until fracture in the stress-strain curves of UFG materials processed by severe plastic deformation (SPD) [27,28] and the values of Y/E for UFG Ti are lower than 0.005 [1], which can explain the obvious pile-up phenomenon for UFG Ti. 3.4. Nanoindentation creep behavior In order to investigate the creep behavior with nanoindentation tests, the data collected during the holding stage was used. The experimental data during the holding stage could be well
Fig. 2. Load-displacement curves for (a) CG Ti, (b) UFG Ti (1P) and (c) UFG Ti (4P).
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Fig. 3. Indentation morphology of (a) CG Ti, (b) UFG Ti (1P) and (c) UFG Ti (4P) at the loading strain rate of 0.1 s
fitted as the following equation [29] 1/2
h = hi + a(t − ti )
1/4
+ b(t − ti )
1/8
+ c (t − ti )
ε̇ =
(1)
where h is the indenter displacement during the holding stage, t is creep time, and hi, ti, a, b and c are the best-fit parameters derived from the Eq. (1). For example, the experimental data and fitted creep displacement and time curve of CG Ti during the nanoindentation holding stage are shown in Fig. 4. Black data points show experimental results while the red line exhibits the fitted curve according to the Eq. (1) and the values of the best-fit parameters are also listed in the Fig. 4. It is clearly seen that the fitted curve shows a good agreement with the experimental results for the ambient-temperature creep of CG Ti in the holding stage. Fig. 5 shows the creep displacement-time (h-t) curves during the nanoindentation creep of CG and UFG Ti at different loading strain rates. It can be observed from Fig. 5 that the creep displacement is the largest for CG Ti and the smallest for UFG Ti after 1 pass of ECAP. It is also noted that the total creep displacement tested at the loading strain rate of 0.005 s 1 is smaller than that tested at the loading strain rate of 0.1 s 1 for CG and UFG Ti. The same results can be obtained from the load plateaus in P-h curves (Fig. 2). For example, at the holding time of 300 s, the largest creep displacement of CG Ti is 70 nm and 110 nm at the loading strain rate of 0.005 s 1 and 0.1 s 1, respectively. Due to high stress level at the higher loading strain rate, lots of dislocations can nucleate, which makes the density of dislocation increase during the loading stage. The dislocation structure formed in the loading stage is highly unstable and tends to relax upon the holding stage in a constant load and leads to significant creep displacements [30]. According to Mayo et al. [31], the creep strain rate ε̇ was calculated from displacement through the following equation
.
(2)
In this work, the creep strain rate data was derived by differentiating the fitted creep h-t curves in the hold stage with Eq. (1) and was plotted in Fig. 6. It can be seen from Fig. 6 that the creep ε̇-t curves consist of transient stage and steady-state stage. At the onset of load holding, the indenter deepens at very high creep strain rate and then the creep strain rate decreases rapidly due to work hardening caused by the plastic deformation, corresponding to the transient creep. After a holding time of 100 s, a steadystate creep stage occurs and is characterized by a very slowly decreased creep strain rate with increasing creep time. The creep strain rate is gradually close to a steady value, which due to the competing mechanisms of work hardening and dynamic recovery [32]. Some researchers have demonstrated that the ambient-temperature creep behavior and dominant creep mechanism of the materials can be analyzed by the simple power-law relationship between the stress (s) and the creep strain rate ( ε̇) [20,33,34]
ε ̇ = ασ n
(3)
where α is a material constant, n is the creep stress exponent. The stress in a nanoindentation test should be related to the pressure exerted by the indenter. From the relationship of H = P /A = P /24.56hc2 (where P is the holding load and A is the projected contact area of the Berkovich tip, which is the crosssectional area of the indenter at the depth to which it is indented, rather than the real surface contact area), a relationship for the creep stress in the holding stage can be easily derived based on σ = (H /3) /(hmax /h)2 [35]. The creep stress exponent is deduced by determining the slope of the ln ε̇ versus lns plot according to Eq. (4).
n=
Fig. 4. Experimental and fitted creep displacement and time curve of CG Ti at a loading strain rate of 0.005 s 1.
ḣ 1 dh = h h dt
1
d(ln ε)̇ d(ln σ )
(4)
Fig. 7 shows the ln ε̇ versus lns plots using the definitions of creep strain rate and stress. In order to obtain the creep stress exponent of CG and UFG Ti, only the steady-state creep is taken into consideration because the value of creep stress exponent differs at different points in the creep strain rate vs. stress curve. The values of the holding load (P), steady creep strain rate ( ε̇) and creep stress exponent (n) at the loading strain rate of 0.005 s 1and 0.1 s 1 are listed in Table 1. In this work, the steady creep strain rates for CG and UFG Ti come out to be (6.94–4.18) 10 5 s 1. It indicates that the steady creep strain rates do not vary significantly with grain size. The insensitivity of the steady creep strain rate with grain size is the signature of dislocation creep [16]. However, the creep stress exponent calculated at the loading strain rate of 0.005 s 1 for CG Ti
X. Liu et al. / Materials Science & Engineering A 676 (2016) 73–79
Fig. 5. Creep displacement-time curves for CG and UFG Ti at different loading strain rates (a) 0.005 s
77
1
; (b) 0.1 s
1
.
Fig. 6. Creep strain rate vs. creep time curves for (a) CG Ti; (b) UFG Ti (1 P); (c) UFG Ti (4 P).
is 9.6 and increases to 18.1 and 20.8 for UFG Ti after 4 passes and 1 pass of ECAP, respectively. The creep stress exponents of CG Ti measured by uniaxial tensile tests increased from 3.0 to 7.0 as the impurity content increased from 0.03 mass% O to 0.09 mass% O and grain size decreased from 100 mm to 17 mm [36,37]. It means that the creep stress exponents of CG Ti are dependent on the impurity O content and grain size. In this work, the 0.06 mass% O was contained in CG Ti with average grain size of about 26 mm. There are differences in impurity O content or grain size with the references. So the creep stress exponents may have a little diffident. Generally, the values of the creep stress exponent calculated by nanoindentation tests are relatively close to those obtained by conventional tensile tests. The creep stress exponent, n is an important parameter because it is commonly used to give indications on the creep stability and dominant creep mechanism. As is commonly known, the creep stress exponent is a value of 1 or 2 at low stresses which corresponding to the diffusion creep
mechanism and grain boundary sliding creep mechanism, respectively. When the dislocation movement dominates the creep deformation, the creep stress exponent is generally larger than 3 [38–40]. It can be seen from the Table 1 that the creep stress exponent varies from 9.6 to 24.6, which indicates that the creep process of CG and UFG Ti is dominated by dislocation creep mechanism. The larger value of the stress exponent the higher creep resistance is. The values of the creep stress exponent of UFG Ti are obviously higher than those of CG Ti, which indicates that UFG Ti has better creep resistance than CG Ti. However, the creep properties of UFG Ti decrease with the number of ECAP passes increasing. The effect of the number of ECAP passes on creep behavior of UFG Ti is consistent with earlier results concerning the creep testing on pure aluminum [11–13] and pure copper [14] processed by ECAP. There is the obvious increase in creep stress exponent with decreasing grain size (from CG to UFG) despite power-law creep is
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Fig. 7. The ln ε̇- lns curves for (a) CG Ti; (b) UFG Ti (1P); (c) UFG Ti (4P). Table 1 The holding load (P), steady creep strain rate (ε̇) and creep stress exponent (n) of CG and UFG Ti. Materials
CG Ti UFG Ti (1P) UFG Ti (4P)
0.005 s
1
0.1 s
P (mN)
ε̇ ( 10
22.09 44.15 45.06
6.94 4.75 5.32
5
s
1
)
1
n
P (mN)
ε̇ ( 10
9.6 20.8 18.1
34.75 49.67 54.89
6.30 4.18 4.97
5
s
1
)
n 13.0 24.6 18.9
expected to be independent of the grain size. It indicates the microstructure has a key role in the creep behavior of Ti and the grain boundaries significantly affect the creep behavior and mechanism. For CG Ti with low dislocation density, the dislocation activation is mainly achieved by new dislocation nucleation and dislocation motion in the grain interiors. The creep strain rates are determined by the grain interiors. New dislocation nucleation and emission becomes the dominant mechanism during plastic deformation [41]. However, when the grain size decreases, the volume of grain boundary zones in the material increases and the grain boundary zone contribution to creep strain rates is higher than the grain interiors. For UFG Ti after 1 pass of ECAP, the initial grains are elongated and refined, and the dislocation density is significantly enhanced. A large number of dislocations were observed at the grain boundaries and the grain interior (Fig. 1b). Kameyanana et al. [37] found that the increased dislocation density by cold-rolling play a major role in suppressing the transient creep and decreasing the steady-state creep strain rate in the ambient-temperature creep of pure Ti. In such a case, the creep process of UFG Ti after 1 pass of ECAP exhibits the lowest steady creep rate of 4.18 10 5 s 1 and the highest creep stress exponent of 24.6 which approximately twice higher than that of CG Ti at the loading strain rate of 0.1 s 1. It is well known that the recovery process
(such as the annihilation of dislocations, and the rearrangement of dislocations at grain boundaries) will occur and the fraction of high angle grain boundaries had obviously increased in UFG Ti during the repetitive pressing [42]. For the UFG Ti after 4 passes of ECAP, recovery process decreases the dislocation density at grain boundaries and grain interior observed by TEM (Fig. 1c), while the grain size is kept at 170 nm. The high angle grain boundaries have a lower strengthening effect on creep resistance because the high angle grain boundaries can be the source to emit and absorb the dislocation [43–45]. Consequently, steady creep strain rate of UFG Ti after 4 pass of ECAP slightly increases and creep stress exponent decreases, compared with the UFG Ti after 1 pass of ECAP. From the Table 1, it can be seen that the creep stress exponents of CG and UFG Ti increase with increasing the loading strain rate, which indicates that the creep stress exponent n exhibits the dependency on the loading strain rate and similar conclusion was proposed by Ma et al. [46] and Li et al. [47]. A possible explanation of this phenomenon is that the corresponding increased holding load with increasing the loading strain rate leads to more defects, such as dislocations. When the effective recovery of the dislocations (such as annihilation and the rearrangement of dislocations) occurs, the creep behavior at larger holding load tends to be difficult to reach the steady state during the holding stage [47].
4. Conclusions The ambient-temperature creep behavior and mechanism of UFG Ti were investigated by nanoindentation test. The results show that an obvious ambient-temperature creep deformation occurs in CG and UFG Ti. The steady creep strain rates do not vary significantly with grain size, while the creep stress exponent
X. Liu et al. / Materials Science & Engineering A 676 (2016) 73–79
values are found to be exceptionally high for UFG Ti, which indicates that UFG Ti processed by ECAP has higher creep resistance than that of CG Ti. The creep behavior of UFG Ti slightly depends on the number of ECAP passes. The creep stress exponent n is proved to increase with increasing the loading strain rate. Based on the values of stress exponent of 18.1–24.6, the ambient-temperature creep mechanism of UFG Ti can be explained with the help of grain boundary mediated dislocation dynamics, especially for dislocation emission and annihilation at grain boundaries.
Acknowledgements The authors acknowledge the financial support from National Nature Science Foundation of China (51474170), Nature Science Foundation of Shaanxi Province of China (2016JQ5026) and Specialized Research Fund of Educational Commission of Shaanxi Province of China (14JK1390).
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Ambient-temperature nanoindentation creep in ultrafine-grained titanium processed by ECAP Xiaoyan Liu n, Qianqian Zhang, Xicheng Zhao, Xirong Yang, Lei Luo School of Metallurgical Engineering, Xi’an University of Architecture and Technology, Xi’an 710055, China
art ic l e i nf o
a b s t r a c t
Article history: Received 12 April 2016 Received in revised form 21 August 2016 Accepted 27 August 2016 Available online 29 August 2016
Ultrafine-grained (UFG) Ti was prepared by equal channel angular pressing (ECAP) through 1 and 4 passes at room temperature. The ambient-temperature creep behavior and mechanism of UFG Ti were studied by nanoindentation creep tests at the loading strain rate of 0.005 and 0.1 s 1. The effect of grain size on creep behavior and relevant creep parameters such as steady creep strain rate and creep stress exponent (n) was estimated for coarse-grained (CG) and UFG Ti. The results show that the creep resistance of UFG Ti is enhanced with respect to CG Ti. UFG Ti after 1 pass of ECAP exhibits the highest creep resistance. The creep resistance of UFG Ti decreases with increasing the number of ECAP passes. The creep stress exponents of UFG Ti are much higher than those of CG Ti while steady creep strain rates do not vary much with grain size. The creep stress exponents of CG and UFG Ti are dependent on the loading strain rate and increase with increasing the loading strain rate. The power-law creep with a stress exponent of 18.1–24.6 is consistent with dislocation process, especially for dislocation emission and annihilation at grain boundaries, which is a key mechanism in creep deformation of UFG Ti. Grain boundaries may play an important role in the creep behavior of UFG Ti. & 2016 Elsevier B.V. All rights reserved.
Keywords: Ultrafine-grained titanium Nanoindentation Creep Dislocations
1. Introduction Titanium and its alloys with high strength and excellent corrosion resistance are widely applied in the fields of aerospace and transport. Ultrafine-grained Ti processed by equal channel angular pressing (ECAP) has been under intensive attention worldwide for past two decades due to its low weight, excellent mechanical properties, and good biocompatibility for potential engineering applications and biomedical implants. To date, a large amount of studies have been reported on the microstructure evolution, deformation behavior and fatigue failure mechanism of UFG Ti [1–3]. Considering perspectives of engineering applications, the creep properties are often of primary importance. Ambient-creep was observed in coarse-grained (CG) titanium about the 1950s [4,5]. Since that time, many studies of ambienttemperature creep of titanium and its alloys have been carried out by materials researchers [6–8]. It was demonstrated that ambienttemperature creep mechanisms of Ti alloys were interpreted as the straightly aligned dislocations slip. Later, Neeraj et al. [9] and Hasija et al. [10] found that ambient-temperature creep behavior of Ti–6Al complied with the Andrade's creep law by experiments and computer simulation, respectively. n
Corresponding author. E-mail address: [email protected] (X. Liu).
http://dx.doi.org/10.1016/j.msea.2016.08.111 0921-5093/& 2016 Elsevier B.V. All rights reserved.
The grain size is expected to play an important role in creep behavior of UFG materials. However, the study on the creep behavior of UFG materials processed by ECAP was scarcely carried out. These reports concerning the creep behavior of ECAP materials were focused on pure aluminum and pure copper [11–14]. It was found that the creep resistance of the UFG materials processed by ECAP was obviously improved with respect to CG materials and the creep behavior of UFG materials strongly depends on the number of ECAP passes [14]. Only very limited results are available describing the creep behavior of UFG Ti processed by ECAP at elevated temperature [15]. The results showed that the high content of grain boundaries and dislocations in the initial state of UFG Ti enhanced its deformation resistance and stabilized the materials. The occurrence of dynamic strain aging in UFG Ti confirmed that the creep mechanism of UFG Ti was related to dislocation motion, such as dislocation slip and climb. Recently, there has been intensive interest to investigate creep behavior and mechanism of micro/nano structural materials and thin films by nanoindentation creep test [16–18]. Although the creep behavior is considerably more complex in nanoindentation creep tests than that in conventional uniaxial tensile/compression tests, it has been found that the creep stress exponent (n) can be calculated from the holding stage in nanoindentation test and the values of stress exponent are generally in good agreement with those obtained by conventional uniaxial tensile/compression tests [19,20]. The purpose of this study is to reveal the ambient-
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temperature creep behavior and mechanism in UFG Ti processed by ECAP using nanoindentation tests.
2. Experimental 2.1. Materials Ultrafine-grained Ti used in this work was prepared by ECAP following route C for 1 and 4 passes at room temperature, respectively. The more details of the processing method were reported elsewhere [21]. Its chemical composition (in wt%) is 0.012 Fe, 0.022 C, 0.003 N, 0.001 H, 0.06 O and balance Ti. The reference coarse-grained (CG) Ti was prepared by annealing the hot rolled pure Ti at 600 °C for 40 min and then air cooling. Microstructure and grain sizes of three materials were examined by transmission electron microscopy (TEM) and optical microscopy (OM). 2.2. Nanoindentation test In order to acquire reliable nanoindentation data, the surface of the indentation specimens was metallographically ground with a
series of progressively finer SiC sandpapers, and mechanically polished to mirror finishing. Nanoindentation tests were performed using the Nano indenter’ XP system (MTS, Inc.) at room temperature. The continuous stiffness measurement (CSM) system was used in this study, which provides a better control on the contact step of the indentation. Upon calibration on standard fused silicon, the tip of the Berkovich diamond indenter was estimated to have a radius of 50 nm. Loads (P) and Displacements (h) were measured with a resolution of 50 nN and 0.01 nm, respectively. All the indents were carried out to a depth limit of 1000 nm at the loading strain rate of 0.005 s 1 and 0.1 s 1, respectively. The indentation direction was parallel to the ECAP axis. In the loading stage, hardness was automatically recorded as a continuous function of surface penetration depth by the CSM mode. After the indenter reached the depth limit, the corresponding load was held for 300 s to monitor the variation in creep displacement with time. Then, the indenter was unloaded to 10% of the maximum load and held for 100 s to correct thermal drift. Finally, the indenter was withdrawn from the specimen surface to terminate the creep test. At least eight indentation tests were carried out at each loading strain rate. Indent surface morphologies were imaged in a ZEISS Supra55 FEG SEM at 15 kV using secondary electron (SE) signals.
Fig. 1. Microstructure of (a) CG Ti, (b) UFG Ti (1P), and (c) UFG Ti (4P).
X. Liu et al. / Materials Science & Engineering A 676 (2016) 73–79
3. Results and discussions 3.1. The microstructures of CG and UFG Ti The microstructures of CG and UFG Ti are shown in Fig. 1. The OM image shows that CG Ti exhibits mostly equiaxed grains with an average grain size of about 26 mm. The micrographs in Figs. 1b and c were taken from the Y plane of ECAP billets. Based on the orthogonal notation used in many earlier studies of ECAP [22], the Y plane is parallel to the side face of the billet at the point of exit from the die. For the UFG Ti after 1 pass of ECAP, elongated parallel shear bands with average widths of about 300 nm are observed in Fig. 1b. Within these parallel bands there are some areas having a high density of dislocations. The TEM image in Fig. 1c shows that the microstructure of UFG Ti after 4 passes of ECAP consists of equiaxed ultrafine grains with high-angle grain boundaries having an average grain size of about 170 nm. Selected area diffraction (SAD) pattern inserted into Fig. 1c also provides evidence for the extent of microstructure refinement. 3.2. Load-displacement curves Fig. 2 shows the representative load-displacement (P-h) curves for CG and UFG Ti at different loading strain rates. In the loading stage, the load required to reach the given displacement increases with increasing the loading strain rate for the CG and UFG Ti. This indicates a pronounced strain rate sensitivity of hardness or stress of three materials. This observation is in agreement with other findings on UFG Ti and Cu [23,24]. At the same loading strain rate, the holding loads of three materials are the highest and lowest for UFG Ti after 4 passes of ECAP and CG Ti, respectively. Furthermore, the hardness values are found to increase with decreasing the grain size, which due to the significant grain refinement and high
75
density of dislocations in UFG Ti after ECAP. Significant load plateaus are observed in P-h curves for CG and UFG Ti. Such load plateau demonstrates a significant creep deformation in the holding stage. The width of the load plateaus increases with increasing the loading strain rate for CG and UFG Ti. 3.3. Indentation morphology Representative SEM indentation micrographs of CG and UFG Ti at the loading strain rate of 0.1 s 1 are shown in Fig. 3. It can be seen that no cracking in any of the indentations, which suggests fully plastic flow in CG and UFG Ti. The indentation morphology of CG Ti (Fig. 3a) shows a fairly uniform surface topography with no visible pile-up. In contrast, the pile-up behavior of UFG Ti is observed in indentation morphology (Fig. 3b and c), indicating inhomogeneous or localized plastic flow. The dimensional analysis and finite element calculations showed that the pile-up and sinkin phenomena were determined by the value of Y/E (where E is the Young's modulus and Y is the initial yield stress), as well as the work-hardening exponent, n [25,26]. Pile-up is significant only when the value of Y/E is very small (Y/E o0.01) and the materials with a small work-hardening exponent (n o0.1) [26]. It is well known that very little work-hardening occurs until fracture in the stress-strain curves of UFG materials processed by severe plastic deformation (SPD) [27,28] and the values of Y/E for UFG Ti are lower than 0.005 [1], which can explain the obvious pile-up phenomenon for UFG Ti. 3.4. Nanoindentation creep behavior In order to investigate the creep behavior with nanoindentation tests, the data collected during the holding stage was used. The experimental data during the holding stage could be well
Fig. 2. Load-displacement curves for (a) CG Ti, (b) UFG Ti (1P) and (c) UFG Ti (4P).
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Fig. 3. Indentation morphology of (a) CG Ti, (b) UFG Ti (1P) and (c) UFG Ti (4P) at the loading strain rate of 0.1 s
fitted as the following equation [29] 1/2
h = hi + a(t − ti )
1/4
+ b(t − ti )
1/8
+ c (t − ti )
ε̇ =
(1)
where h is the indenter displacement during the holding stage, t is creep time, and hi, ti, a, b and c are the best-fit parameters derived from the Eq. (1). For example, the experimental data and fitted creep displacement and time curve of CG Ti during the nanoindentation holding stage are shown in Fig. 4. Black data points show experimental results while the red line exhibits the fitted curve according to the Eq. (1) and the values of the best-fit parameters are also listed in the Fig. 4. It is clearly seen that the fitted curve shows a good agreement with the experimental results for the ambient-temperature creep of CG Ti in the holding stage. Fig. 5 shows the creep displacement-time (h-t) curves during the nanoindentation creep of CG and UFG Ti at different loading strain rates. It can be observed from Fig. 5 that the creep displacement is the largest for CG Ti and the smallest for UFG Ti after 1 pass of ECAP. It is also noted that the total creep displacement tested at the loading strain rate of 0.005 s 1 is smaller than that tested at the loading strain rate of 0.1 s 1 for CG and UFG Ti. The same results can be obtained from the load plateaus in P-h curves (Fig. 2). For example, at the holding time of 300 s, the largest creep displacement of CG Ti is 70 nm and 110 nm at the loading strain rate of 0.005 s 1 and 0.1 s 1, respectively. Due to high stress level at the higher loading strain rate, lots of dislocations can nucleate, which makes the density of dislocation increase during the loading stage. The dislocation structure formed in the loading stage is highly unstable and tends to relax upon the holding stage in a constant load and leads to significant creep displacements [30]. According to Mayo et al. [31], the creep strain rate ε̇ was calculated from displacement through the following equation
.
(2)
In this work, the creep strain rate data was derived by differentiating the fitted creep h-t curves in the hold stage with Eq. (1) and was plotted in Fig. 6. It can be seen from Fig. 6 that the creep ε̇-t curves consist of transient stage and steady-state stage. At the onset of load holding, the indenter deepens at very high creep strain rate and then the creep strain rate decreases rapidly due to work hardening caused by the plastic deformation, corresponding to the transient creep. After a holding time of 100 s, a steadystate creep stage occurs and is characterized by a very slowly decreased creep strain rate with increasing creep time. The creep strain rate is gradually close to a steady value, which due to the competing mechanisms of work hardening and dynamic recovery [32]. Some researchers have demonstrated that the ambient-temperature creep behavior and dominant creep mechanism of the materials can be analyzed by the simple power-law relationship between the stress (s) and the creep strain rate ( ε̇) [20,33,34]
ε ̇ = ασ n
(3)
where α is a material constant, n is the creep stress exponent. The stress in a nanoindentation test should be related to the pressure exerted by the indenter. From the relationship of H = P /A = P /24.56hc2 (where P is the holding load and A is the projected contact area of the Berkovich tip, which is the crosssectional area of the indenter at the depth to which it is indented, rather than the real surface contact area), a relationship for the creep stress in the holding stage can be easily derived based on σ = (H /3) /(hmax /h)2 [35]. The creep stress exponent is deduced by determining the slope of the ln ε̇ versus lns plot according to Eq. (4).
n=
Fig. 4. Experimental and fitted creep displacement and time curve of CG Ti at a loading strain rate of 0.005 s 1.
ḣ 1 dh = h h dt
1
d(ln ε)̇ d(ln σ )
(4)
Fig. 7 shows the ln ε̇ versus lns plots using the definitions of creep strain rate and stress. In order to obtain the creep stress exponent of CG and UFG Ti, only the steady-state creep is taken into consideration because the value of creep stress exponent differs at different points in the creep strain rate vs. stress curve. The values of the holding load (P), steady creep strain rate ( ε̇) and creep stress exponent (n) at the loading strain rate of 0.005 s 1and 0.1 s 1 are listed in Table 1. In this work, the steady creep strain rates for CG and UFG Ti come out to be (6.94–4.18) 10 5 s 1. It indicates that the steady creep strain rates do not vary significantly with grain size. The insensitivity of the steady creep strain rate with grain size is the signature of dislocation creep [16]. However, the creep stress exponent calculated at the loading strain rate of 0.005 s 1 for CG Ti
X. Liu et al. / Materials Science & Engineering A 676 (2016) 73–79
Fig. 5. Creep displacement-time curves for CG and UFG Ti at different loading strain rates (a) 0.005 s
77
1
; (b) 0.1 s
1
.
Fig. 6. Creep strain rate vs. creep time curves for (a) CG Ti; (b) UFG Ti (1 P); (c) UFG Ti (4 P).
is 9.6 and increases to 18.1 and 20.8 for UFG Ti after 4 passes and 1 pass of ECAP, respectively. The creep stress exponents of CG Ti measured by uniaxial tensile tests increased from 3.0 to 7.0 as the impurity content increased from 0.03 mass% O to 0.09 mass% O and grain size decreased from 100 mm to 17 mm [36,37]. It means that the creep stress exponents of CG Ti are dependent on the impurity O content and grain size. In this work, the 0.06 mass% O was contained in CG Ti with average grain size of about 26 mm. There are differences in impurity O content or grain size with the references. So the creep stress exponents may have a little diffident. Generally, the values of the creep stress exponent calculated by nanoindentation tests are relatively close to those obtained by conventional tensile tests. The creep stress exponent, n is an important parameter because it is commonly used to give indications on the creep stability and dominant creep mechanism. As is commonly known, the creep stress exponent is a value of 1 or 2 at low stresses which corresponding to the diffusion creep
mechanism and grain boundary sliding creep mechanism, respectively. When the dislocation movement dominates the creep deformation, the creep stress exponent is generally larger than 3 [38–40]. It can be seen from the Table 1 that the creep stress exponent varies from 9.6 to 24.6, which indicates that the creep process of CG and UFG Ti is dominated by dislocation creep mechanism. The larger value of the stress exponent the higher creep resistance is. The values of the creep stress exponent of UFG Ti are obviously higher than those of CG Ti, which indicates that UFG Ti has better creep resistance than CG Ti. However, the creep properties of UFG Ti decrease with the number of ECAP passes increasing. The effect of the number of ECAP passes on creep behavior of UFG Ti is consistent with earlier results concerning the creep testing on pure aluminum [11–13] and pure copper [14] processed by ECAP. There is the obvious increase in creep stress exponent with decreasing grain size (from CG to UFG) despite power-law creep is
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Fig. 7. The ln ε̇- lns curves for (a) CG Ti; (b) UFG Ti (1P); (c) UFG Ti (4P). Table 1 The holding load (P), steady creep strain rate (ε̇) and creep stress exponent (n) of CG and UFG Ti. Materials
CG Ti UFG Ti (1P) UFG Ti (4P)
0.005 s
1
0.1 s
P (mN)
ε̇ ( 10
22.09 44.15 45.06
6.94 4.75 5.32
5
s
1
)
1
n
P (mN)
ε̇ ( 10
9.6 20.8 18.1
34.75 49.67 54.89
6.30 4.18 4.97
5
s
1
)
n 13.0 24.6 18.9
expected to be independent of the grain size. It indicates the microstructure has a key role in the creep behavior of Ti and the grain boundaries significantly affect the creep behavior and mechanism. For CG Ti with low dislocation density, the dislocation activation is mainly achieved by new dislocation nucleation and dislocation motion in the grain interiors. The creep strain rates are determined by the grain interiors. New dislocation nucleation and emission becomes the dominant mechanism during plastic deformation [41]. However, when the grain size decreases, the volume of grain boundary zones in the material increases and the grain boundary zone contribution to creep strain rates is higher than the grain interiors. For UFG Ti after 1 pass of ECAP, the initial grains are elongated and refined, and the dislocation density is significantly enhanced. A large number of dislocations were observed at the grain boundaries and the grain interior (Fig. 1b). Kameyanana et al. [37] found that the increased dislocation density by cold-rolling play a major role in suppressing the transient creep and decreasing the steady-state creep strain rate in the ambient-temperature creep of pure Ti. In such a case, the creep process of UFG Ti after 1 pass of ECAP exhibits the lowest steady creep rate of 4.18 10 5 s 1 and the highest creep stress exponent of 24.6 which approximately twice higher than that of CG Ti at the loading strain rate of 0.1 s 1. It is well known that the recovery process
(such as the annihilation of dislocations, and the rearrangement of dislocations at grain boundaries) will occur and the fraction of high angle grain boundaries had obviously increased in UFG Ti during the repetitive pressing [42]. For the UFG Ti after 4 passes of ECAP, recovery process decreases the dislocation density at grain boundaries and grain interior observed by TEM (Fig. 1c), while the grain size is kept at 170 nm. The high angle grain boundaries have a lower strengthening effect on creep resistance because the high angle grain boundaries can be the source to emit and absorb the dislocation [43–45]. Consequently, steady creep strain rate of UFG Ti after 4 pass of ECAP slightly increases and creep stress exponent decreases, compared with the UFG Ti after 1 pass of ECAP. From the Table 1, it can be seen that the creep stress exponents of CG and UFG Ti increase with increasing the loading strain rate, which indicates that the creep stress exponent n exhibits the dependency on the loading strain rate and similar conclusion was proposed by Ma et al. [46] and Li et al. [47]. A possible explanation of this phenomenon is that the corresponding increased holding load with increasing the loading strain rate leads to more defects, such as dislocations. When the effective recovery of the dislocations (such as annihilation and the rearrangement of dislocations) occurs, the creep behavior at larger holding load tends to be difficult to reach the steady state during the holding stage [47].
4. Conclusions The ambient-temperature creep behavior and mechanism of UFG Ti were investigated by nanoindentation test. The results show that an obvious ambient-temperature creep deformation occurs in CG and UFG Ti. The steady creep strain rates do not vary significantly with grain size, while the creep stress exponent
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values are found to be exceptionally high for UFG Ti, which indicates that UFG Ti processed by ECAP has higher creep resistance than that of CG Ti. The creep behavior of UFG Ti slightly depends on the number of ECAP passes. The creep stress exponent n is proved to increase with increasing the loading strain rate. Based on the values of stress exponent of 18.1–24.6, the ambient-temperature creep mechanism of UFG Ti can be explained with the help of grain boundary mediated dislocation dynamics, especially for dislocation emission and annihilation at grain boundaries.
Acknowledgements The authors acknowledge the financial support from National Nature Science Foundation of China (51474170), Nature Science Foundation of Shaanxi Province of China (2016JQ5026) and Specialized Research Fund of Educational Commission of Shaanxi Province of China (14JK1390).
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