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Vacancy-assisted hardening in nanostructured metals
Materials Letters 65 (2011) 514–516
Contents lists available at ScienceDirect
Materials Letters j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m a t l e t
Vacancy-assisted hardening in nanostructured metals L.H. Su a,b, C. Lu a,⁎, A.K. Tieu a, L.Z. He b, Y. Zhang c, D. Wexler a a b c
School of Mechanical, Materials and Mechatronic Engineering, University of Wollongong, Wollongong, NSW 2522, Australia School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China Institute for Superconducting and Electronic Materials, University of Wollongong, Fairy Meadow 2519, Australia
a r t i c l e
i n f o
Article history: Received 5 October 2010 Accepted 31 October 2010 Available online 4 November 2010 Keywords: Nanostructured materials Severe plastic deformation Point defects Deformation mechanism
a b s t r a c t Bulk metals deformed by severe plastic deformation have some unusual phenomenon such as hardening by low temperature annealing. Commercial purity aluminum processed by equal channel angular pressing has been used to investigate the deformation mechanism during further annealing of the severed plastic deformed materials. It has been found that vacancy-type defects induced by severe plastic deformation are responsible for some unusual phenomena, such as hardening by annealing and softening by slight deformation in nanostructured metals. It is anticipated that the effective control of vacancy-type defects during SPD and consequent annealing process will be combined to attain an optimum combination of strength and ductility in nanostructured metals. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Bulk nanostructured metals with grain sizes of 100–500 nm, processed by severe plastic deformation, have attracted a great deal of attention in the last decade since they exhibit great mechanical properties, such as significantly enhanced strength and hardness over those of their coarse-grained counterparts [1–3]. Special interests have been paid on the hardening mechanism of nanostructured metals. It is commonly accepted that annealing leads to a decrease in strength. An unusual annealing-induced hardening has been observed in nanostructured metals [4–7]. However, the hardening mechanism is still not clear. In the current research, commercial purity aluminum processed by equal channel angular pressing was examined by X-ray diffraction and differential scanning calorimetry so as to investigate the hardening mechanism during further annealing of the severe plastic deformed materials. Understanding of this hardening mechanism will gain a deeper insight of nanostructured metals and will also provide explanation to other special phenomena observed in the nanostructured metals. 2. Experimental The current experiments were conducted using commercial purity aluminum 1050. The material was fully annealed at 723 K for 1 h and cut into 20× 20× 100 mm3 billets. The billets were then pressed through an
⁎ Corresponding author. Faculty of Engineering, University of Wollongong, Northfield Ave, Wollongong, NSW 2522, Australia. Tel.: + 61 2 42214639. E-mail address: [email protected] (C. Lu). 0167-577X/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2010.10.090
ECAP die with the intersection angle 90° and the outer arc of curvature 20°. The pressings were continued for seven passes via route BC to obtain an equivalent strain of 7. The six- and seven-pass samples were annealed at 80 °C, 110 °C, 150 °C and 180 °C for 4 h. Tensile specimens with a gauge dimension of 8 × 3 × 1.5 mm3 were cut from the as-ECAPed and further annealed materials with the tensile axis parallel to the pressing direction. Tensile tests were conducted on an Instron 5566 testing machine with an initial strain rate of 5 × 10− 4 s− 1. Specimens for XRD were cut from the central part of the billets in the same direction as tensile specimens. The specimens were ground and electropolished before subjected to XRD tests. The XRD patterns were obtained via a GBC MMA X-ray diffractometer. DSC analysis was performed in METTLER TOELDO TGA/DSC-1 thermogravimetric analyser protected with flowing Argon. 3. Results and discussion Yield stresses (0.2% proof stresses) are plotted against annealing temperature in Fig. 1. It can be seen that for the seven-pass sample hardening occurs at 80 °C and 110 °C and softening at 150 °C and 180 °C. The six-pass sample exhibits hardening at 80 °C, 110 °C and 150 °C and softening at 180 °C. The maximum yield stress for the six- and sevenpass sample occurs at 110 °C and 80 °C, respectively. This indicates that high strain shifts the peak annealing temperature to a lower value. XRD was conducted on the ECAPed samples and Rietveld refinement was performed to analyse the lattice constants. The refined lattice constants are 4.0497 Å, 4.0493 Å, 4.0491 Å, 4.0485 Å, 4.0473 Å and 4.0465 Å for 0, 1, 2, 4, 6 and 7 passes ECAPed samples, respectively. It is known that the lattice constant is not sensitive to the line defects or planar defects, but directly related to the existence of
L.H. Su et al. / Materials Letters 65 (2011) 514–516
515
Fig. 1. Yield stresses of six and seven-pass ECAP samples annealed at different temperatures.
point defects [8]. In cubic crystals, the lattice constant and concentrations of point defects are related by: 3ðΔa = aÞ = αv cv + αi ci
ð1Þ
where a is the lattice constant and Δa is its change, αv and αi are the relaxation volumes for vacancies and interstitials in concentrations cv and ci, respectively [8]. The decrease in lattice constants indicates the existence of vacancies. Vacancy-type defects have been observed in severely deformed and thin films aluminum alloys [9,10]. When two dislocations on adjacent slip planes meet each other during plastic deformation, a string of vacancies or self-interstitials is generated. Since the formation energy for self-interstitials (2.59 eV) is approximately four times of that for vacancies (0.68 eV) [11], it is hard to form self-interstitials in aluminum. Even if self-interstitials can be generated, they will easily annihilate in the sinks because the activation energy for migration of self-interstitials (0.1 eV [8]) is much smaller than that of vacancies (0.64 eV [11]). Therefore, the concentration of self-interstitials is expected to be negligible. It has been reported that in commercial purity aluminum, nanoparticles can precipitate in the grain boundaries and prevent annihilation of vacancies [12]. The generated vacancies will then be accumulated as the deformation proceeds, result in super-saturated vacancies in highly strained samples. Super-saturated vacancy type defects have also been found in nanostructured Ni and Cu [13,14]. Fig. 2 shows the DSC curves of seven-pass sample at heating rates 3, 5 and 8 °C/min, respectively. Two exothermic peaks appear in Fig. 2(A) and they significantly overlap in Fig. 2(B) and (C). Gaussian fittings with two peaks, named by P1 and P2, were performed, and are shown by dash lines. The peaks P1 and P2 are centered at about 100–122 °C and 142–188 °C, which have been identified to the annealing of vacancies and dislocations, respectively [13,15]. Compared to the yield stresses in Fig. 1, the peak P1 can be associated to hardening around 80–110 °C, and the peak P2 to softening at 150 °C and above. Fig. 3 shows the activation energy Q as the function of the energy dependent transformed fraction Y for the peaks P1 and P2. It can be seen that at the early stage of the peak P1 Q increases from 0.47 eV at Y = 10% to 0.51 eV at Y = 20%, and then decreases with Y. The trend is similar with Ref. [15], in which the migration of vacancies has been identified to be responsible for the reaction. The activation energy for the early stage peak P1 is close to, but smaller than the migrating energy of mono-vacancy (0.57–0.62 eV) [15]. This implies that vacancy migration dominates at the early stage of the peak P1 but it is not the unique reaction. There should be another exothermal reaction responsible for the decrease of activation energy. After the vacancies are activated, randomly moving vacancies can meet both dislocations and other vacancies. If the vacancies diffuse to the core of
Fig. 2. DSC curves of seven-pass ECAP sample at different heating rates: (A) 3 °C/min, (B) 5 °C/min, and (C) 8 °C/min.
a dislocation the dislocation climbs and will finally annihilate in the sinks, leading to a reduced dislocation density. If a mono-vacancy meets another mono-vacancy, the smallest vacancy cluster, di-vacancy, will be created. If a vacancy meets an existing vacancy cluster, the cluster will grow to become the stacking fault tetrahedral (SFT), vacancy Frank loop, or even void. SFT and vacancy Frank loop are more stable [8,16]. They may not be mobile at the temperature range of the peak P1. The formation energy of a stationary vacancy cluster includes two parts: migration energy of the mono-vacancies and binding energy between the mono-vacancies and the existing vacancy clusters. For a stationary di-vacancy, the formation energy (EF2) can be expressed by: EF2 = EM1 –EB2
ð2Þ
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L.H. Su et al. / Materials Letters 65 (2011) 514–516
annihilate at the sinks, which leads to the reduction of the number of dislocation pinning centers and hence softening of the materials.
4. Conclusions
Fig. 3. Activation energy Q as functions of the energy dependent transformed fraction Y.
where EM1 is the migration energy of a mono-vacancy and EB2 is the binding energy between two mono-vacancies. The equilibrium concentration measurement gave EM1 = 0.62 eV and EB2 = 0.25 eV, which yield EF2 = 0.37 eV [17]. As the vacancy cluster grows up, the formation energy EF will decrease due to the continuous binding. Therefore, the activation energy decreases at the middle and final stages of the peak P1. This conclusion is consistent with Ref. [18], which states that at small concentrations of vacancies the annealing process is controlled by the migration energy of single vacancies whereas at large densities of vacancies the annealing process is controlled by the migration energy of vacancies and binding energy to form vacancy clusters. The stationary large vacancy clusters will act as pinning centers and hinder the slip motion of the dislocations, results in enhanced strength. Therefore, hardening by annealing at the low temperature can be attributed to the dislocation pinning effect caused by the formed vacancy clusters. When the annealing temperature increases to the early stage of the peak P2, the vacancy clusters will be thermally activated. As can be seen in Fig. 3, the activation energy for the peak P2 is much smaller than the migration energy of the mono-vacancy at the early stage. This is consistent with Matsukawa and Zinkle's observation that the effective migration energy for the glissile vacancy cluster is much less than that for mono-vacancy [19]. The glissile vacancy clusters will
Aforementioned analysis indicates that severe plastic deformation can induce super-saturated vacancies in metals. During the low temperature annealing the thermally activated vacancies trigger the dislocation climb and the vacancy clusters grow. The former leads to the reduction of dislocation density whereas the latter results in hardening by dislocation pinning effect, which is the hardening by low temperature annealing effect. Deeper understanding of the influence of vacancy-type defects on nanostructured metals may inspire the development of new processing methods to achieve superior mechanical properties.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
Valiev RZ, Langdon TG. Prog Mater Sci 2006;51:881–981. Zhu YT, Langdon TG. JOM 2004;56:58–63. Tsuji N, Saito Y, Lee SH, Minamono Y. Adv Eng Mater 2003;53:338–44. Horita Z, Fujinami T, Nemoto M, Langdon TG. Metall Mater Trans A 2000;31: 691–701. Huang X, Hansen N, Tsuji N. Science 2006;312:249–51. Bowen JR. Mater Sci Eng A 2008;483–484:231–4. Kamikawa N, Huang X, Tsuji N, Hansen N. Acta Mater 2009;57:4198–208. Damask AC, Dienes GJ. Point defects in metals. New York: Gordon and Beach, Science Publishers; 1971. p. 170–2. Wu XL, Li B, Ma E. Appl Phys Lett 2007;91:141908. Ohkubo H, Shimomura Y, Mukouda I, Sugio K, Kiritani M. Mater Sci Eng A 2003;350:30–6. Mishin Y, Farkas D, Mehl MJ, Papaconstantopoulos DA. Phys Rev B 1999;59: 3393–407. Saimoto S, Cooley J, Larsen H, Scholler C. Philos Mag 2009;89:853–68. Setman D, Schafler E, Korznikova E, Zehetbauer MJ. Mater Sci Eng A 2008;493: 116–22. Ungár T, Schafler E, Hanák P, Bernstorff S, Zehetbauer M. Mater Sci Eng A 2007;462:398–401. Schmidt J, Haeβner F. Z Phys B Con Mat 1990;81:215–22. Gavini V, Bhattacharya K, Ortiz M. Phys Rev B 2007;76:180101. Seeger A, Schumacher D, Schilling W, Diehl J. Vacancies and interstitials in metals. Amsterdam: North Holland Pub Co; 1969. p. 38–42. Seeger A, Schumacher D, Schilling W, Diehl J. Vacancies and interstitials in metals. Amsterdam: North Holland Pub Co; 1969. p. 215–49. Matsukawa Y, Zinkle SJ. Science 2007;318:959–62.
Contents lists available at ScienceDirect
Materials Letters j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m a t l e t
Vacancy-assisted hardening in nanostructured metals L.H. Su a,b, C. Lu a,⁎, A.K. Tieu a, L.Z. He b, Y. Zhang c, D. Wexler a a b c
School of Mechanical, Materials and Mechatronic Engineering, University of Wollongong, Wollongong, NSW 2522, Australia School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China Institute for Superconducting and Electronic Materials, University of Wollongong, Fairy Meadow 2519, Australia
a r t i c l e
i n f o
Article history: Received 5 October 2010 Accepted 31 October 2010 Available online 4 November 2010 Keywords: Nanostructured materials Severe plastic deformation Point defects Deformation mechanism
a b s t r a c t Bulk metals deformed by severe plastic deformation have some unusual phenomenon such as hardening by low temperature annealing. Commercial purity aluminum processed by equal channel angular pressing has been used to investigate the deformation mechanism during further annealing of the severed plastic deformed materials. It has been found that vacancy-type defects induced by severe plastic deformation are responsible for some unusual phenomena, such as hardening by annealing and softening by slight deformation in nanostructured metals. It is anticipated that the effective control of vacancy-type defects during SPD and consequent annealing process will be combined to attain an optimum combination of strength and ductility in nanostructured metals. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Bulk nanostructured metals with grain sizes of 100–500 nm, processed by severe plastic deformation, have attracted a great deal of attention in the last decade since they exhibit great mechanical properties, such as significantly enhanced strength and hardness over those of their coarse-grained counterparts [1–3]. Special interests have been paid on the hardening mechanism of nanostructured metals. It is commonly accepted that annealing leads to a decrease in strength. An unusual annealing-induced hardening has been observed in nanostructured metals [4–7]. However, the hardening mechanism is still not clear. In the current research, commercial purity aluminum processed by equal channel angular pressing was examined by X-ray diffraction and differential scanning calorimetry so as to investigate the hardening mechanism during further annealing of the severe plastic deformed materials. Understanding of this hardening mechanism will gain a deeper insight of nanostructured metals and will also provide explanation to other special phenomena observed in the nanostructured metals. 2. Experimental The current experiments were conducted using commercial purity aluminum 1050. The material was fully annealed at 723 K for 1 h and cut into 20× 20× 100 mm3 billets. The billets were then pressed through an
⁎ Corresponding author. Faculty of Engineering, University of Wollongong, Northfield Ave, Wollongong, NSW 2522, Australia. Tel.: + 61 2 42214639. E-mail address: [email protected] (C. Lu). 0167-577X/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2010.10.090
ECAP die with the intersection angle 90° and the outer arc of curvature 20°. The pressings were continued for seven passes via route BC to obtain an equivalent strain of 7. The six- and seven-pass samples were annealed at 80 °C, 110 °C, 150 °C and 180 °C for 4 h. Tensile specimens with a gauge dimension of 8 × 3 × 1.5 mm3 were cut from the as-ECAPed and further annealed materials with the tensile axis parallel to the pressing direction. Tensile tests were conducted on an Instron 5566 testing machine with an initial strain rate of 5 × 10− 4 s− 1. Specimens for XRD were cut from the central part of the billets in the same direction as tensile specimens. The specimens were ground and electropolished before subjected to XRD tests. The XRD patterns were obtained via a GBC MMA X-ray diffractometer. DSC analysis was performed in METTLER TOELDO TGA/DSC-1 thermogravimetric analyser protected with flowing Argon. 3. Results and discussion Yield stresses (0.2% proof stresses) are plotted against annealing temperature in Fig. 1. It can be seen that for the seven-pass sample hardening occurs at 80 °C and 110 °C and softening at 150 °C and 180 °C. The six-pass sample exhibits hardening at 80 °C, 110 °C and 150 °C and softening at 180 °C. The maximum yield stress for the six- and sevenpass sample occurs at 110 °C and 80 °C, respectively. This indicates that high strain shifts the peak annealing temperature to a lower value. XRD was conducted on the ECAPed samples and Rietveld refinement was performed to analyse the lattice constants. The refined lattice constants are 4.0497 Å, 4.0493 Å, 4.0491 Å, 4.0485 Å, 4.0473 Å and 4.0465 Å for 0, 1, 2, 4, 6 and 7 passes ECAPed samples, respectively. It is known that the lattice constant is not sensitive to the line defects or planar defects, but directly related to the existence of
L.H. Su et al. / Materials Letters 65 (2011) 514–516
515
Fig. 1. Yield stresses of six and seven-pass ECAP samples annealed at different temperatures.
point defects [8]. In cubic crystals, the lattice constant and concentrations of point defects are related by: 3ðΔa = aÞ = αv cv + αi ci
ð1Þ
where a is the lattice constant and Δa is its change, αv and αi are the relaxation volumes for vacancies and interstitials in concentrations cv and ci, respectively [8]. The decrease in lattice constants indicates the existence of vacancies. Vacancy-type defects have been observed in severely deformed and thin films aluminum alloys [9,10]. When two dislocations on adjacent slip planes meet each other during plastic deformation, a string of vacancies or self-interstitials is generated. Since the formation energy for self-interstitials (2.59 eV) is approximately four times of that for vacancies (0.68 eV) [11], it is hard to form self-interstitials in aluminum. Even if self-interstitials can be generated, they will easily annihilate in the sinks because the activation energy for migration of self-interstitials (0.1 eV [8]) is much smaller than that of vacancies (0.64 eV [11]). Therefore, the concentration of self-interstitials is expected to be negligible. It has been reported that in commercial purity aluminum, nanoparticles can precipitate in the grain boundaries and prevent annihilation of vacancies [12]. The generated vacancies will then be accumulated as the deformation proceeds, result in super-saturated vacancies in highly strained samples. Super-saturated vacancy type defects have also been found in nanostructured Ni and Cu [13,14]. Fig. 2 shows the DSC curves of seven-pass sample at heating rates 3, 5 and 8 °C/min, respectively. Two exothermic peaks appear in Fig. 2(A) and they significantly overlap in Fig. 2(B) and (C). Gaussian fittings with two peaks, named by P1 and P2, were performed, and are shown by dash lines. The peaks P1 and P2 are centered at about 100–122 °C and 142–188 °C, which have been identified to the annealing of vacancies and dislocations, respectively [13,15]. Compared to the yield stresses in Fig. 1, the peak P1 can be associated to hardening around 80–110 °C, and the peak P2 to softening at 150 °C and above. Fig. 3 shows the activation energy Q as the function of the energy dependent transformed fraction Y for the peaks P1 and P2. It can be seen that at the early stage of the peak P1 Q increases from 0.47 eV at Y = 10% to 0.51 eV at Y = 20%, and then decreases with Y. The trend is similar with Ref. [15], in which the migration of vacancies has been identified to be responsible for the reaction. The activation energy for the early stage peak P1 is close to, but smaller than the migrating energy of mono-vacancy (0.57–0.62 eV) [15]. This implies that vacancy migration dominates at the early stage of the peak P1 but it is not the unique reaction. There should be another exothermal reaction responsible for the decrease of activation energy. After the vacancies are activated, randomly moving vacancies can meet both dislocations and other vacancies. If the vacancies diffuse to the core of
Fig. 2. DSC curves of seven-pass ECAP sample at different heating rates: (A) 3 °C/min, (B) 5 °C/min, and (C) 8 °C/min.
a dislocation the dislocation climbs and will finally annihilate in the sinks, leading to a reduced dislocation density. If a mono-vacancy meets another mono-vacancy, the smallest vacancy cluster, di-vacancy, will be created. If a vacancy meets an existing vacancy cluster, the cluster will grow to become the stacking fault tetrahedral (SFT), vacancy Frank loop, or even void. SFT and vacancy Frank loop are more stable [8,16]. They may not be mobile at the temperature range of the peak P1. The formation energy of a stationary vacancy cluster includes two parts: migration energy of the mono-vacancies and binding energy between the mono-vacancies and the existing vacancy clusters. For a stationary di-vacancy, the formation energy (EF2) can be expressed by: EF2 = EM1 –EB2
ð2Þ
516
L.H. Su et al. / Materials Letters 65 (2011) 514–516
annihilate at the sinks, which leads to the reduction of the number of dislocation pinning centers and hence softening of the materials.
4. Conclusions
Fig. 3. Activation energy Q as functions of the energy dependent transformed fraction Y.
where EM1 is the migration energy of a mono-vacancy and EB2 is the binding energy between two mono-vacancies. The equilibrium concentration measurement gave EM1 = 0.62 eV and EB2 = 0.25 eV, which yield EF2 = 0.37 eV [17]. As the vacancy cluster grows up, the formation energy EF will decrease due to the continuous binding. Therefore, the activation energy decreases at the middle and final stages of the peak P1. This conclusion is consistent with Ref. [18], which states that at small concentrations of vacancies the annealing process is controlled by the migration energy of single vacancies whereas at large densities of vacancies the annealing process is controlled by the migration energy of vacancies and binding energy to form vacancy clusters. The stationary large vacancy clusters will act as pinning centers and hinder the slip motion of the dislocations, results in enhanced strength. Therefore, hardening by annealing at the low temperature can be attributed to the dislocation pinning effect caused by the formed vacancy clusters. When the annealing temperature increases to the early stage of the peak P2, the vacancy clusters will be thermally activated. As can be seen in Fig. 3, the activation energy for the peak P2 is much smaller than the migration energy of the mono-vacancy at the early stage. This is consistent with Matsukawa and Zinkle's observation that the effective migration energy for the glissile vacancy cluster is much less than that for mono-vacancy [19]. The glissile vacancy clusters will
Aforementioned analysis indicates that severe plastic deformation can induce super-saturated vacancies in metals. During the low temperature annealing the thermally activated vacancies trigger the dislocation climb and the vacancy clusters grow. The former leads to the reduction of dislocation density whereas the latter results in hardening by dislocation pinning effect, which is the hardening by low temperature annealing effect. Deeper understanding of the influence of vacancy-type defects on nanostructured metals may inspire the development of new processing methods to achieve superior mechanical properties.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
Valiev RZ, Langdon TG. Prog Mater Sci 2006;51:881–981. Zhu YT, Langdon TG. JOM 2004;56:58–63. Tsuji N, Saito Y, Lee SH, Minamono Y. Adv Eng Mater 2003;53:338–44. Horita Z, Fujinami T, Nemoto M, Langdon TG. Metall Mater Trans A 2000;31: 691–701. Huang X, Hansen N, Tsuji N. Science 2006;312:249–51. Bowen JR. Mater Sci Eng A 2008;483–484:231–4. Kamikawa N, Huang X, Tsuji N, Hansen N. Acta Mater 2009;57:4198–208. Damask AC, Dienes GJ. Point defects in metals. New York: Gordon and Beach, Science Publishers; 1971. p. 170–2. Wu XL, Li B, Ma E. Appl Phys Lett 2007;91:141908. Ohkubo H, Shimomura Y, Mukouda I, Sugio K, Kiritani M. Mater Sci Eng A 2003;350:30–6. Mishin Y, Farkas D, Mehl MJ, Papaconstantopoulos DA. Phys Rev B 1999;59: 3393–407. Saimoto S, Cooley J, Larsen H, Scholler C. Philos Mag 2009;89:853–68. Setman D, Schafler E, Korznikova E, Zehetbauer MJ. Mater Sci Eng A 2008;493: 116–22. Ungár T, Schafler E, Hanák P, Bernstorff S, Zehetbauer M. Mater Sci Eng A 2007;462:398–401. Schmidt J, Haeβner F. Z Phys B Con Mat 1990;81:215–22. Gavini V, Bhattacharya K, Ortiz M. Phys Rev B 2007;76:180101. Seeger A, Schumacher D, Schilling W, Diehl J. Vacancies and interstitials in metals. Amsterdam: North Holland Pub Co; 1969. p. 38–42. Seeger A, Schumacher D, Schilling W, Diehl J. Vacancies and interstitials in metals. Amsterdam: North Holland Pub Co; 1969. p. 215–49. Matsukawa Y, Zinkle SJ. Science 2007;318:959–62.