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Influence of the recovery and recrystallization processes on the martensitic transformation of cold worked equiatomic Ti–Ni alloy
Materials Science and Engineering A355 (2003) 292 /298 www.elsevier.com/locate/msea
Influence of the recovery and recrystallization processes on the martensitic transformation of cold worked equiatomic Ti Ni alloy /
F. Khelfaoui 1, G. Gue´nin * GEMPPM Laboratory, UMR 5510, INSA-Lyon, Batiment Blaise Pascal, 7 Avenue Jean Capelle, 69621 Villeurbanne, Cedex, France Received 24 May 2002; received in revised form 22 January 2003
Abstract The martensitic transformation of Ti /Ni shape memory alloys is very sensitive to thermomechanical treatments (cold work and annealings). In the present paper, a Ti /Ni alloy close to the equiatomic composition has been 40% cold rolled and then submitted to various annealing treatments, each one characterized by temperature and time (Ta, ta). A large range of annealing temperatures (538 /778 K) and times (10 s /123 d) has been covered. From differential scanning calorimetry measurements, three zones have been identified with reference to (Ta, ta) values. Zone I, at low temperatures and times, is characterized by badly defined transformations on cooling as well as on heating; it corresponds to a recovery /reverse martensitic transformation stage. Zone II exhibits large changes of the transformation features which occur in two steps on cooling (R phase then martensitic) and in one step on heating; this zone corresponds to the recrystallization and growth of very small stressed grains. Similar transformation behavior can be obtained with different (Ta, ta) annealings, which can be characterized by an activation energy of :/3.4 eV in the range 698 /778 K. Zone III relates to a well defined martensitic transformation taking place in one step with almost no further evolution; it corresponds to the growth of stress-free grains. # 2003 Elsevier Science B.V. All rights reserved. Keywords: Ti /Ni shape memory alloy; Thermomechanical treatments; Recovery; Recrystallization; R-phase; Martensitic transformation
1. Introduction For various commercial applications, the Ti/Ni shape memory alloys are often cold worked [1]. This process introduces defects, such as dislocations, that for large strain amplitudes, inhibits the martensitic transformation. During the subsequent thermal treatments, the martensitic transformation is able to regenerate. As already shown in the literature, this occurs with some changes, in particular the emergence of a two step transformation on cooling for low temperature annealings: austenite 0/R phase 0/martensite [2 /5]. Much research has been carried out to study the effect of * Corresponding author. Tel.: /33-472-438-245; fax: /33-472-427930. E-mail address: [email protected] (G. Gue´nin). 1 Currently at Max-Planck-Institut fu¨ r Metallforschung, Heisenbergstraße 1, D-70569 Stuttgart, Germany.
heat treatments on the evolution of the martensitic transformation [5 /8]. However, no systematic study has been performed in order to explain the changes of the martensitic transformation during the recovery and recrystallization processes. According to the microstructure evolution followed by thermoelectric power (TEP), two heat treatment stages have been defined elsewhere by the present authors [9]. These domains have been identified as corresponding to recovery and recrystallization processes with special features. The first domain is observed for low annealing temperatures (T B/650 K) and is characterized by an apparent activation energy of :/2.7 eV. The second domain is observed for higher annealing temperatures (T /650 K) and is characterized by an apparent activation energy of :/3.6 eV. The aim of the present paper is to observe the evolution of the martensitic transformation properties during these annealing treatments.
0921-5093/03/$ - see front matter # 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0921-5093(03)00068-6
F. Khelfaoui, G. Gue´nin / Materials Science and Engineering A355 (2003) 292 /298
2. Experimental procedure A Ti/49.60 at.% Ni alloy, provided by the CEA (Atomic Energy Center) of Grenoble, is used for this study. Plates 66 /4 /1 mm3 are first annealed for 2 h at 1143 K, quenched in water and then 40% cold rolled in the martensitic state at room temperature. Small samples of 15 mg weight, used for differential scanning calorimetry (DSC), are carefully cut from this plate with a low speed diamond saw. The remaining plate is used for thermoelectric power (TEP) measurements, as described in Ref. [9] and microhardness HV300. Both samples are annealed in a temperature range from 538 to 778 K with 40 K intervals, for cumulated times ranging from 10 s to 123 d. The samples are then slightly polished to remove the oxide layer before each measurement. The DSC measurements are performed from 200 to 395 K, with heating and cooling rates of 5 K min 1, using a Mettler 3000 apparatus. The HV hardness measurements are carried out in the martensite state, with a WOLPERT HV-TESTOR machine using a load of 300 g, during 15 s. Each value is the average of five measurements.
3. Results 3.1. DSC measurements and behavior of the martensite transformation Samples as annealed (2 h at 1143 K) exhibit only one transformation, as shown in Fig. 1(a). On cooling, the
Fig. 1. DSC thermograms of Ti /49.60 at.% Ni alloy: (a) annealed 2 h at 1143 K; (b) after 40% room temperature cold rolling.
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austenite to martensite transformation peak is noted (AM) and on heating, the martensite to austenite transformation peak is noted (MA). Just after cold work, no clear transformation is observed in the temperature range explored (Fig. 1b). Fig. 2 shows typical DSC curves observed following various annealing treatments (more results are presented in Ref. [10]). After only 1 min at 778 K, the sample clearly exhibits the martensitic transformation (Fig. 2a). On cooling, two well separated peaks appear corresponding to the two transformations austenite 0/R phase (AR) and R phase 0/martensite (RM), respectively. The (AR) peak is narrow and well defined and its temperature does not change with the annealing time. On the contrary, the (RM) peak is broad, it becomes sharper and its temperature increases when increasing the annealing time. On heating, a single peak is present associated to the martensite 0/austenite transformation (MA). It becomes sharper and its temperature shows little changes when increasing the annealing time. The temperature difference between the (AR) and (RM) peaks decreases with the annealing time until the two peaks become superimposed after an estimated annealing time of :/4 h at 778 K. For longer annealing times, the classic (AM) peak is observed and its temperature remains stable. For a 698 K annealing temperature (Fig. 2b), the martensitic transformation behavior is similar to the one observed at 778 K, but with longer annealing times. The lower annealing temperature allows for a better observation of the beginning of the martensitic transformation evolution. At the beginning, the reverse transformation (MA) exhibits a broad peak. During cooling, the (AR) peak is small but rather sharp, whereas the (RM) peak is strongly flattened. The three peaks (MA), (AR) and (RM) exhibit the same evolution as previously described for a 778 K annealing treatment. After a long time (123 d), the (AR) and (RM) peaks merge together to form a single peak (AM). For the 738 K annealing treatment, which is not presented in this paper, the transformation behavior is intermediate between those for 698 and 778 K annealing treatments and the merging of the (AR) peak and the (RM) peak occurs after 37 h. The thermograms obtained after 621 and 538 K annealing temperatures are shown in Fig. 2(c, d, respectively). These thermograms are typical of the zone corresponding to low annealing temperatures. It can be seen that for short annealing times, the transformations are very little regenerated, whereas for longer annealing times, an evolution begins which is comparable to the one obtained for higher annealing temperatures (698 and 778 K). By continuity it is possible to identify the (AR), (RM) and (MA) peaks, even though they are very diffuse. From these results, three annealing zones can be defined as a function of (Ta, ta) temperature and time of annealings:
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Fig. 2. Evolution of DSC thermograms of Ti /49.60 at.% Ni alloy after 40% cold rolling at room temperature and annealing treatments at: (a) 778 K; (b) 698 K; (c) 621 K; (d) 538 K.
/ Zone I: the (RM) peak is badly defined. The (AR) and (MA) peaks are better observed and increase progressively, however their heights stay at a relatively low level if compared to fully grown peaks. / Zone II: the (RM) peak can be identified even if it is broad and the (AR) and (MA) peaks are well defined. In this zone, the main change is a progressive sharpening of all the peaks and a shift of the (RM) peak towards the high temperatures with an increase of the annealing time at a given annealing temperature. On the contrary, the two other peaks do not suffer significant temperature changes. / Zone III: the (RM) peak has merged with the (AR) peak giving a unique (AM) peak on cooling, therefore
only the (AM) and (MA) peaks are observed with no further noticeable evolution. These zones are shown in Fig. 3. The plots, corresponding to the transition recovery-recrystallization deduced from thermoelectric power measurements, are also shown in Fig. 3 [9]. There is a rather good agreement between these plots and the transition from zones I to II. In zone II, it seems that there is a comparable evolution of the thermograms for different annealing temperatures, but with different time scales: for example an annealing of 1 min at 778 K is nearly equivalent to a treatment of 4 h at 698 K. This can be quantified if the
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Fig. 2 (Continued)
(RM) peak temperature is plotted as a function of time, for different annealing temperatures, as shown in Fig. 4. For each annealing temperature, the increase of (RM) temperature is followed by a saturation which corresponds to the merging of the (RM) peak with the (AR) peak (arrows in Fig. 4). These data can be analyzed with the empirical Johnson /Mehl /Avrami formalism [18], as in the previous work for the TEP data [9]. The fraction change follows the general expression: y 1 exp[ (kta )n ]; in the present case, y is defined as the relative increase of the (RM) peak temperature (Fig. 4), k depends only on the annealing temperature Ta, ta is the annealing time and n is a constant. A reasonable fit of this law is obtained which gives an n value of 0.199/ 0.02 for all the annealing temperatures. The Arrhenius plots ln k vs. 1/Ta (or ta vs. 1/Ta for different fixed values of y ) are not well aligned for the whole annealing temperature domain. However, if only the annealing temperatures /658 K are taken into account, an
apparent activation energy of 3.49/0.4 eV is found. This value and the corresponding n value (0.199/0.02) are in fair agreement with the values 3.59/0.3 and 0.199/ 0.02 eV associated with the changes of the TEP in the same temperature range [9]. In Fig. 5, the TEP values as a function of the (RM) peak temperature are plotted; it can be seen that a correlation is present in zone II, especially for the three highest annealing temperatures ranging from 698 to 778 K, which confirms the upper result. 3.2. Microhardness measurements Fig. 6 illustrates the evolution of the microhardness HV300 following the different annealing temperatures as a function of the annealing time. A decrease of HV300 with increasing annealing times is first observed then, for annealing temperatures /658 K, a plateau at 180 HV is attained. The transition points between the
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Fig. 3. Temperature-time diagram showing the three annealing zones for Ti /49.60 at.% Ni alloy after 40% room temperature cold rolling. Zone I: badly defined transformations; zone II: well defined (AR) and (MA) transformations; large temperature changes of the (MR) transformation peak; zone III: only well defined (AM) and (MA) transformations. Circles correspond to the transition between recovery and recrystallization as deduced from TEP measurements [9].
Fig. 5. Correlation between TEP and (RM) peak temperature measurements for different annealing temperatures: 2 621 K, \ 658 K, ' 698 K, m 738 K, j 778 K. A better correlation is seen for higher annealing temperatures (full symbols).
Fig. 6. Changes of microhardness HV300 of the Ti /49.60 at.% Ni alloy as a function of annealing time after 40% cold rolling for different annealing temperatures: / 539 K, / 578 K, 2 621 K, \ 658 K, ^ 698 K, k 738 K, I 778 K. Arrows indicate the (AR) peak elimination time. The lines are drawn as a guide for the eyes. Fig. 4. Changes of (RM) peak temperature of Ti /49.60 at.% Ni alloy as a function of annealing time after 40% cold rolling for different annealing temperatures: 2 621 K, \ 658 K, ^ 698 K, k 738 K, I 778 K. Arrows indicate the (AR) peak elimination time. The lines are drawn as a guide for the eyes.
decrease of microhardness and the plateau are indicated by arrows in Fig. 5 for 778, 738 and 698 K annealing temperatures. The curves are shifted towards longer times when the temperature is lowered and qualitatively two regimes are observed roughly below and above 621 K; they correspond to apparent activation energies of 2.29/0.4 and 1.49/0.2 eV, respectively.
4. Discussion It is well known that severe cold work, such as 40%, inhibits the martensitic transformation by the introduction of defects, which are essentially dislocations.
Following cold work, the annealing treatment, as defined by temperature and time (Ta, ta), induces different features of the martensitic transformation related to the recovery and recrystallization processes. In zone I (Fig. 3), small recovery of the martensitic transformation takes place and little quantitative information can be deduced. In Fig. 4, no information at all concerning the (RM) peak temperature is found for 538 and 578 K annealing temperatures in this zone whatever the time; the values given for 621 K and to a less extent for 658 K, suffer a large uncertainty. This behavior is in good agreement with those of Lin et al. [11]. Concerning the hardness changes, the curves for 538, 578 and 621 K are in this zone I (apart from the three or four last points for 621 K). The corresponding activation energy of 2.29/0.4 eV is not far from the value 2.89/0.2 eV found by thermoelectric power in the same zone for the same alloy [9], which has been identified as a recovery /reverse martensitic transformation process.
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In zone II, as stated earlier, it appears that similar behaviors in terms of transformation temperatures are encountered for different (Ta, ta) annealing parameters (Fig. 4). Moreover, for Ta within the range 698 /778 K, there is a rather good correlation with the TEP measurements. This could indicate that these annealing treatments correspond to similar features of the microstructure. However, some differences are noticed; for example the temperature width of the (RM) peak, when observed at the same temperature, becomes larger when it is obtained for lower annealing temperature (see, e.g. the peaks for 1 min at 778 K and the peak for 4 h at 698 K). Moreover, in this zone, the microhardness behaves very differently if compared with the transformation behavior (the apparent activation energies are 1.4 and 3.4 eV, respectively). This implies different microhardnesses for similar transformation behaviors obtained by different annealing temperatures. If the same previous example is considered, it can be seen that the hardness is 260 HV for 1 min at 778 K and only 230 HV for 4 h at 698 K. This clearly indicates that for these ‘equivalent treatments’, even if some features are common giving similar transformation behavior and similar TEP, the microstructures must exhibit some differences. It has been shown [9] that in this zone II, at 778 K, the microstructure is formed of very small grains with a Table 1 Synthesis of the results
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diameter increasing from 50 up to 500 nm with time; these grains exhibit some internal stresses. Motohashi et al. [12] have studied the effect of the grain size on the Ms temperature for the Ti/50.2 at.% Ni alloy; these different grain sizes are obtained for 998 and 1273 K annealing temperatures, during different times. They show that, for grain sizes ranging from 4 to 15 mm, Ms decreases of B/20 K, whereas for grain sizes ranging from 15 to 50 mm, Ms increases of about the same value. In the case of the smaller grains for Motohashi et al., the behavior is therefore the opposite of that observed in the present work in zone II. However, it is clear that the annealing temperatures as well as the corresponding grain sizes are not in the same range: the smallest grains in Ref. [12] correspond to 13 d at 778 K, which is clearly situated in zone III of well annealed stress-free grains. This is an indication that the internal stresses observed inside the very small grains, rather than their size, are responsible for an opposite behavior. Following this hypothesis, a possible explanation of the very different behaviors between the kinetics of the transformation changes and of the microhardnesses could be due to the fact that the martensitic transformation (and TEP) should be essentially sensitive to internal stresses, whereas the microhardness should be sensitive to internal stresses and also to grain size (Hall /Petch law).
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The transition from zone II to zone III (two peaks (AR)/(RM) changing into one (AM) peak) corresponds roughly on the one hand, to the disappearance of the internal stresses in recrystallized grains (verified in Ref. [9] for 778 K annealing) and on the other hand, to the final term of the decrease of microhardness. Several studies have reported that the appearance of the R phase could be due to dislocations [3,4,7,13/17]. The R phase seems to nucleate preferentially in the stress concentrated regions [16]. The fact that the (AR) transformation disappears approximately when the internal stresses disappear from the grains seems to corroborate this hypothesis. However, the present results also show that the (AR) transformation is inhibited by a strong defect density. Indeed, the (AR) transformation does not occur just after cold work but after low temperature annealings (538 and 578 K). This has also been previously observed by Larnicol [15]. A synthesis of all the previous observations is given in Table 1.
5. Conclusion This work, in addition to the previous work [9], presents an exhaustive study of the effect of annealing treatments after plastic deformation of Ti/Ni alloys. Three annealing zones characterized by (Ta, ta) couples have been evidenced: / Zone I of badly defined displacive transformations that corresponds to a combined recovery /reverse martensitic transformation stage. / Zone II of large changes of the martensitic transformation features which corresponds to the crystallization and growth of very small stressed grains. Similar transformation behaviors can be obtained above 658 K with different (Ta, ta) annealings, which can be characterized by an activation energy of :/3.4 eV. However, these similar behaviors do not correspond to similar hardnesses, which implies different
features in the microstructure. Some more work is needed on this particular point. / Zone III in which a well defined martensitic transformation occurs in one step with almost no noticeable further evolution and which corresponds to the growth of stress free grains.
References [1] F. Funakubo, Shape Memory Alloys, Gordon and Breach, London, 1987, p. 273. [2] H.C. Lin, S.K. Wu, Metall. Trans. A24 (1993) 293. [3] P.E. Thoma, M.-Y. Kao, S. Fariabi, D.N. Abujudom, in: C.M. Wayman, J. Perkins (Eds.), Proceedings of the International Conference on Martensitic Transformations, Monterey Institute of Advanced Studies, USA, 1992, p. 917. [4] S.K. Wu, H.C. Lin, Y.C. Yen, Mater. Sci. Eng. A215 (1996) 113. [5] P.E. Thoma, D.R. Angst, K.D. Schachner, J. Phys. IV 5 (C8) (1995) 557. [6] D.N. Abujudom, P.E. Thoma, S. Fariabi, Mater. Sci. For. 56 /58 (1990) 565. [7] S.K. Wu, H.C. Lin, P.C. Cheng, J. Mater. Sci. 34 (1999) 5669. [8] C. Somsen, H. Za¨hres, J. Ka¨stner, E.F. Wassermann, T. Kakeshita, T. Saburi, Mater. Sci. Eng. A273 (1999) 310. [9] F. Khelfaoui, G. Thollet, G. Gue´nin, Mater. Sci. Eng. A338 (2002) 297. [10] F. Khelfaoui, Ph.D. Thesis. Institut National des Sciences Applique´es de Lyon, Villeurbanne, France, 2000, 142 pp. [11] H.C. Lin, S.K. Wu, T.S. Chou, H.P. Kao, Acta Metall. Mater. 39 (1991) 2069. [12] Y. Motohashi, T. Sakuma, M. Suzuki, T. Hoshiaya, K. Ohsawa, in: C.M. Wayman, J. Perkins (Eds.), Proceedings of the International Conference on Martensitic Transformations, Monterey Institute of Advanced Studies, USA, 1992, p. 993. [13] Y. Liu, Ph.D. Thesis. University of Western Australia, Nedlands, Australia, 1990, 250 pp. [14] E. Hornbogen, E. Kobus, Z. Metallkd. 87 (1996) 442. [15] M. Larnicol, Ph.D. Thesis. Universite´ Paris XI Orsay, Paris, France, 1998, 202 pp. [16] T. Fukuda, T. Saburi, K. Dio, S. Nenno, Mater. Trans. JIM 33 (1992) 271. [17] H. Matsumoto, J. Mater. Sci. Lett. 10 (1990) 408. [18] J. Burke, La Cine´tique des Changements de Phase dans les Me´taux, Masson, Paris, 1968, p. 247.
Influence of the recovery and recrystallization processes on the martensitic transformation of cold worked equiatomic Ti Ni alloy /
F. Khelfaoui 1, G. Gue´nin * GEMPPM Laboratory, UMR 5510, INSA-Lyon, Batiment Blaise Pascal, 7 Avenue Jean Capelle, 69621 Villeurbanne, Cedex, France Received 24 May 2002; received in revised form 22 January 2003
Abstract The martensitic transformation of Ti /Ni shape memory alloys is very sensitive to thermomechanical treatments (cold work and annealings). In the present paper, a Ti /Ni alloy close to the equiatomic composition has been 40% cold rolled and then submitted to various annealing treatments, each one characterized by temperature and time (Ta, ta). A large range of annealing temperatures (538 /778 K) and times (10 s /123 d) has been covered. From differential scanning calorimetry measurements, three zones have been identified with reference to (Ta, ta) values. Zone I, at low temperatures and times, is characterized by badly defined transformations on cooling as well as on heating; it corresponds to a recovery /reverse martensitic transformation stage. Zone II exhibits large changes of the transformation features which occur in two steps on cooling (R phase then martensitic) and in one step on heating; this zone corresponds to the recrystallization and growth of very small stressed grains. Similar transformation behavior can be obtained with different (Ta, ta) annealings, which can be characterized by an activation energy of :/3.4 eV in the range 698 /778 K. Zone III relates to a well defined martensitic transformation taking place in one step with almost no further evolution; it corresponds to the growth of stress-free grains. # 2003 Elsevier Science B.V. All rights reserved. Keywords: Ti /Ni shape memory alloy; Thermomechanical treatments; Recovery; Recrystallization; R-phase; Martensitic transformation
1. Introduction For various commercial applications, the Ti/Ni shape memory alloys are often cold worked [1]. This process introduces defects, such as dislocations, that for large strain amplitudes, inhibits the martensitic transformation. During the subsequent thermal treatments, the martensitic transformation is able to regenerate. As already shown in the literature, this occurs with some changes, in particular the emergence of a two step transformation on cooling for low temperature annealings: austenite 0/R phase 0/martensite [2 /5]. Much research has been carried out to study the effect of * Corresponding author. Tel.: /33-472-438-245; fax: /33-472-427930. E-mail address: [email protected] (G. Gue´nin). 1 Currently at Max-Planck-Institut fu¨ r Metallforschung, Heisenbergstraße 1, D-70569 Stuttgart, Germany.
heat treatments on the evolution of the martensitic transformation [5 /8]. However, no systematic study has been performed in order to explain the changes of the martensitic transformation during the recovery and recrystallization processes. According to the microstructure evolution followed by thermoelectric power (TEP), two heat treatment stages have been defined elsewhere by the present authors [9]. These domains have been identified as corresponding to recovery and recrystallization processes with special features. The first domain is observed for low annealing temperatures (T B/650 K) and is characterized by an apparent activation energy of :/2.7 eV. The second domain is observed for higher annealing temperatures (T /650 K) and is characterized by an apparent activation energy of :/3.6 eV. The aim of the present paper is to observe the evolution of the martensitic transformation properties during these annealing treatments.
0921-5093/03/$ - see front matter # 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0921-5093(03)00068-6
F. Khelfaoui, G. Gue´nin / Materials Science and Engineering A355 (2003) 292 /298
2. Experimental procedure A Ti/49.60 at.% Ni alloy, provided by the CEA (Atomic Energy Center) of Grenoble, is used for this study. Plates 66 /4 /1 mm3 are first annealed for 2 h at 1143 K, quenched in water and then 40% cold rolled in the martensitic state at room temperature. Small samples of 15 mg weight, used for differential scanning calorimetry (DSC), are carefully cut from this plate with a low speed diamond saw. The remaining plate is used for thermoelectric power (TEP) measurements, as described in Ref. [9] and microhardness HV300. Both samples are annealed in a temperature range from 538 to 778 K with 40 K intervals, for cumulated times ranging from 10 s to 123 d. The samples are then slightly polished to remove the oxide layer before each measurement. The DSC measurements are performed from 200 to 395 K, with heating and cooling rates of 5 K min 1, using a Mettler 3000 apparatus. The HV hardness measurements are carried out in the martensite state, with a WOLPERT HV-TESTOR machine using a load of 300 g, during 15 s. Each value is the average of five measurements.
3. Results 3.1. DSC measurements and behavior of the martensite transformation Samples as annealed (2 h at 1143 K) exhibit only one transformation, as shown in Fig. 1(a). On cooling, the
Fig. 1. DSC thermograms of Ti /49.60 at.% Ni alloy: (a) annealed 2 h at 1143 K; (b) after 40% room temperature cold rolling.
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austenite to martensite transformation peak is noted (AM) and on heating, the martensite to austenite transformation peak is noted (MA). Just after cold work, no clear transformation is observed in the temperature range explored (Fig. 1b). Fig. 2 shows typical DSC curves observed following various annealing treatments (more results are presented in Ref. [10]). After only 1 min at 778 K, the sample clearly exhibits the martensitic transformation (Fig. 2a). On cooling, two well separated peaks appear corresponding to the two transformations austenite 0/R phase (AR) and R phase 0/martensite (RM), respectively. The (AR) peak is narrow and well defined and its temperature does not change with the annealing time. On the contrary, the (RM) peak is broad, it becomes sharper and its temperature increases when increasing the annealing time. On heating, a single peak is present associated to the martensite 0/austenite transformation (MA). It becomes sharper and its temperature shows little changes when increasing the annealing time. The temperature difference between the (AR) and (RM) peaks decreases with the annealing time until the two peaks become superimposed after an estimated annealing time of :/4 h at 778 K. For longer annealing times, the classic (AM) peak is observed and its temperature remains stable. For a 698 K annealing temperature (Fig. 2b), the martensitic transformation behavior is similar to the one observed at 778 K, but with longer annealing times. The lower annealing temperature allows for a better observation of the beginning of the martensitic transformation evolution. At the beginning, the reverse transformation (MA) exhibits a broad peak. During cooling, the (AR) peak is small but rather sharp, whereas the (RM) peak is strongly flattened. The three peaks (MA), (AR) and (RM) exhibit the same evolution as previously described for a 778 K annealing treatment. After a long time (123 d), the (AR) and (RM) peaks merge together to form a single peak (AM). For the 738 K annealing treatment, which is not presented in this paper, the transformation behavior is intermediate between those for 698 and 778 K annealing treatments and the merging of the (AR) peak and the (RM) peak occurs after 37 h. The thermograms obtained after 621 and 538 K annealing temperatures are shown in Fig. 2(c, d, respectively). These thermograms are typical of the zone corresponding to low annealing temperatures. It can be seen that for short annealing times, the transformations are very little regenerated, whereas for longer annealing times, an evolution begins which is comparable to the one obtained for higher annealing temperatures (698 and 778 K). By continuity it is possible to identify the (AR), (RM) and (MA) peaks, even though they are very diffuse. From these results, three annealing zones can be defined as a function of (Ta, ta) temperature and time of annealings:
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Fig. 2. Evolution of DSC thermograms of Ti /49.60 at.% Ni alloy after 40% cold rolling at room temperature and annealing treatments at: (a) 778 K; (b) 698 K; (c) 621 K; (d) 538 K.
/ Zone I: the (RM) peak is badly defined. The (AR) and (MA) peaks are better observed and increase progressively, however their heights stay at a relatively low level if compared to fully grown peaks. / Zone II: the (RM) peak can be identified even if it is broad and the (AR) and (MA) peaks are well defined. In this zone, the main change is a progressive sharpening of all the peaks and a shift of the (RM) peak towards the high temperatures with an increase of the annealing time at a given annealing temperature. On the contrary, the two other peaks do not suffer significant temperature changes. / Zone III: the (RM) peak has merged with the (AR) peak giving a unique (AM) peak on cooling, therefore
only the (AM) and (MA) peaks are observed with no further noticeable evolution. These zones are shown in Fig. 3. The plots, corresponding to the transition recovery-recrystallization deduced from thermoelectric power measurements, are also shown in Fig. 3 [9]. There is a rather good agreement between these plots and the transition from zones I to II. In zone II, it seems that there is a comparable evolution of the thermograms for different annealing temperatures, but with different time scales: for example an annealing of 1 min at 778 K is nearly equivalent to a treatment of 4 h at 698 K. This can be quantified if the
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Fig. 2 (Continued)
(RM) peak temperature is plotted as a function of time, for different annealing temperatures, as shown in Fig. 4. For each annealing temperature, the increase of (RM) temperature is followed by a saturation which corresponds to the merging of the (RM) peak with the (AR) peak (arrows in Fig. 4). These data can be analyzed with the empirical Johnson /Mehl /Avrami formalism [18], as in the previous work for the TEP data [9]. The fraction change follows the general expression: y 1 exp[ (kta )n ]; in the present case, y is defined as the relative increase of the (RM) peak temperature (Fig. 4), k depends only on the annealing temperature Ta, ta is the annealing time and n is a constant. A reasonable fit of this law is obtained which gives an n value of 0.199/ 0.02 for all the annealing temperatures. The Arrhenius plots ln k vs. 1/Ta (or ta vs. 1/Ta for different fixed values of y ) are not well aligned for the whole annealing temperature domain. However, if only the annealing temperatures /658 K are taken into account, an
apparent activation energy of 3.49/0.4 eV is found. This value and the corresponding n value (0.199/0.02) are in fair agreement with the values 3.59/0.3 and 0.199/ 0.02 eV associated with the changes of the TEP in the same temperature range [9]. In Fig. 5, the TEP values as a function of the (RM) peak temperature are plotted; it can be seen that a correlation is present in zone II, especially for the three highest annealing temperatures ranging from 698 to 778 K, which confirms the upper result. 3.2. Microhardness measurements Fig. 6 illustrates the evolution of the microhardness HV300 following the different annealing temperatures as a function of the annealing time. A decrease of HV300 with increasing annealing times is first observed then, for annealing temperatures /658 K, a plateau at 180 HV is attained. The transition points between the
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Fig. 3. Temperature-time diagram showing the three annealing zones for Ti /49.60 at.% Ni alloy after 40% room temperature cold rolling. Zone I: badly defined transformations; zone II: well defined (AR) and (MA) transformations; large temperature changes of the (MR) transformation peak; zone III: only well defined (AM) and (MA) transformations. Circles correspond to the transition between recovery and recrystallization as deduced from TEP measurements [9].
Fig. 5. Correlation between TEP and (RM) peak temperature measurements for different annealing temperatures: 2 621 K, \ 658 K, ' 698 K, m 738 K, j 778 K. A better correlation is seen for higher annealing temperatures (full symbols).
Fig. 6. Changes of microhardness HV300 of the Ti /49.60 at.% Ni alloy as a function of annealing time after 40% cold rolling for different annealing temperatures: / 539 K, / 578 K, 2 621 K, \ 658 K, ^ 698 K, k 738 K, I 778 K. Arrows indicate the (AR) peak elimination time. The lines are drawn as a guide for the eyes. Fig. 4. Changes of (RM) peak temperature of Ti /49.60 at.% Ni alloy as a function of annealing time after 40% cold rolling for different annealing temperatures: 2 621 K, \ 658 K, ^ 698 K, k 738 K, I 778 K. Arrows indicate the (AR) peak elimination time. The lines are drawn as a guide for the eyes.
decrease of microhardness and the plateau are indicated by arrows in Fig. 5 for 778, 738 and 698 K annealing temperatures. The curves are shifted towards longer times when the temperature is lowered and qualitatively two regimes are observed roughly below and above 621 K; they correspond to apparent activation energies of 2.29/0.4 and 1.49/0.2 eV, respectively.
4. Discussion It is well known that severe cold work, such as 40%, inhibits the martensitic transformation by the introduction of defects, which are essentially dislocations.
Following cold work, the annealing treatment, as defined by temperature and time (Ta, ta), induces different features of the martensitic transformation related to the recovery and recrystallization processes. In zone I (Fig. 3), small recovery of the martensitic transformation takes place and little quantitative information can be deduced. In Fig. 4, no information at all concerning the (RM) peak temperature is found for 538 and 578 K annealing temperatures in this zone whatever the time; the values given for 621 K and to a less extent for 658 K, suffer a large uncertainty. This behavior is in good agreement with those of Lin et al. [11]. Concerning the hardness changes, the curves for 538, 578 and 621 K are in this zone I (apart from the three or four last points for 621 K). The corresponding activation energy of 2.29/0.4 eV is not far from the value 2.89/0.2 eV found by thermoelectric power in the same zone for the same alloy [9], which has been identified as a recovery /reverse martensitic transformation process.
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In zone II, as stated earlier, it appears that similar behaviors in terms of transformation temperatures are encountered for different (Ta, ta) annealing parameters (Fig. 4). Moreover, for Ta within the range 698 /778 K, there is a rather good correlation with the TEP measurements. This could indicate that these annealing treatments correspond to similar features of the microstructure. However, some differences are noticed; for example the temperature width of the (RM) peak, when observed at the same temperature, becomes larger when it is obtained for lower annealing temperature (see, e.g. the peaks for 1 min at 778 K and the peak for 4 h at 698 K). Moreover, in this zone, the microhardness behaves very differently if compared with the transformation behavior (the apparent activation energies are 1.4 and 3.4 eV, respectively). This implies different microhardnesses for similar transformation behaviors obtained by different annealing temperatures. If the same previous example is considered, it can be seen that the hardness is 260 HV for 1 min at 778 K and only 230 HV for 4 h at 698 K. This clearly indicates that for these ‘equivalent treatments’, even if some features are common giving similar transformation behavior and similar TEP, the microstructures must exhibit some differences. It has been shown [9] that in this zone II, at 778 K, the microstructure is formed of very small grains with a Table 1 Synthesis of the results
297
diameter increasing from 50 up to 500 nm with time; these grains exhibit some internal stresses. Motohashi et al. [12] have studied the effect of the grain size on the Ms temperature for the Ti/50.2 at.% Ni alloy; these different grain sizes are obtained for 998 and 1273 K annealing temperatures, during different times. They show that, for grain sizes ranging from 4 to 15 mm, Ms decreases of B/20 K, whereas for grain sizes ranging from 15 to 50 mm, Ms increases of about the same value. In the case of the smaller grains for Motohashi et al., the behavior is therefore the opposite of that observed in the present work in zone II. However, it is clear that the annealing temperatures as well as the corresponding grain sizes are not in the same range: the smallest grains in Ref. [12] correspond to 13 d at 778 K, which is clearly situated in zone III of well annealed stress-free grains. This is an indication that the internal stresses observed inside the very small grains, rather than their size, are responsible for an opposite behavior. Following this hypothesis, a possible explanation of the very different behaviors between the kinetics of the transformation changes and of the microhardnesses could be due to the fact that the martensitic transformation (and TEP) should be essentially sensitive to internal stresses, whereas the microhardness should be sensitive to internal stresses and also to grain size (Hall /Petch law).
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The transition from zone II to zone III (two peaks (AR)/(RM) changing into one (AM) peak) corresponds roughly on the one hand, to the disappearance of the internal stresses in recrystallized grains (verified in Ref. [9] for 778 K annealing) and on the other hand, to the final term of the decrease of microhardness. Several studies have reported that the appearance of the R phase could be due to dislocations [3,4,7,13/17]. The R phase seems to nucleate preferentially in the stress concentrated regions [16]. The fact that the (AR) transformation disappears approximately when the internal stresses disappear from the grains seems to corroborate this hypothesis. However, the present results also show that the (AR) transformation is inhibited by a strong defect density. Indeed, the (AR) transformation does not occur just after cold work but after low temperature annealings (538 and 578 K). This has also been previously observed by Larnicol [15]. A synthesis of all the previous observations is given in Table 1.
5. Conclusion This work, in addition to the previous work [9], presents an exhaustive study of the effect of annealing treatments after plastic deformation of Ti/Ni alloys. Three annealing zones characterized by (Ta, ta) couples have been evidenced: / Zone I of badly defined displacive transformations that corresponds to a combined recovery /reverse martensitic transformation stage. / Zone II of large changes of the martensitic transformation features which corresponds to the crystallization and growth of very small stressed grains. Similar transformation behaviors can be obtained above 658 K with different (Ta, ta) annealings, which can be characterized by an activation energy of :/3.4 eV. However, these similar behaviors do not correspond to similar hardnesses, which implies different
features in the microstructure. Some more work is needed on this particular point. / Zone III in which a well defined martensitic transformation occurs in one step with almost no noticeable further evolution and which corresponds to the growth of stress free grains.
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