《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Martensitic phase transformations in nanocrystalline NiTi

Availableonlineatwww.sciencedirect.com BCIENCE Acta materialia ELSEVIER Acta Materialia 52(2004)137-147 www.actamat-journals.com Martensitic phase transformations in nanocrystalline Niti studied by tem T. Waitz.v Kazykhanov. H P. Karnthaler Received 16 June 2003: accepted 22 August 2003 Abstract By high pressure torsion(HPT) deformation almost complete amorphization is obtained in bulk Ni-50. 3at %Ti containing B1 martensite. During low temperature annealing tiny crystallites retained after the hpt deformation are acting as nuclei and trigger the nanocrystallization of B2 austenite. It is shown that the density of the nuclei is a function of the HPt strain and determines together with the annealing temperature the grain size of the nanocrystals ranging from 5 to 350 nm. Upon cooling the nano- structures transform to b19 partially since the grain boundaries hinder the autocatalytic formation of martensite. The large transformation strains of B19 are reduced by very fine(00 1) compound twins. With decreasing grain size an increasing energy barrier arises and the martensitic transformation is completely suppressed in grains smaller than 60 nm. The r-phase transformation causing only small transformation strains is observed in grains between 15 and 60 nm. Whereas in grains below 15 nm B2 remains dicating that no transformation occurs at all e 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: High pressure torsion; Bulk amorphous materials; Nanocrystallization; Martensitic phase transformation: High-resolution electron microscopy 1. ntroduction cause of the nanocrystallization in the HPt deformed alloys remained unclear. In the intermetallic compound NiTi the shape mem- As the martensitic transformation occurring in ory effect and superelastic properties are related to a nanocrystalline NiTi is concerned there are contradic martensitic phase transformation being of considerable tory results presented in the literature. In Ni-498at %Ti interest both from a scientific and a technological point subjected to cold rolling crystal refinement and partial of view. In addition, in NiTi amorphization can be in- amorphization was achieved and the martensitic trans duced applying various solid state processes such as formation was suppressed even after cooling to-150C particle irradiation [1] and strong mechanical deforma- [5, 6]. The reason for the observed change in the trans tion by cold rolling [2] and mechanical alloying [3]. formation seems to be unclear; it was proposed that Recently severe plastic deformation by high pressure dislocations induced by the cold rolling may stabilize the torsion(HPT) methods was applied to NiTi achieving parent B2 austenite. Contrary to this, it was reported amorphization; followed by a suitable heat treatment a at the martensite start temperature Ms increases with nanocrystalline phase can obtained in bulk HPT alloys decreasing grain size in a bulk nanocrystalline material by devitrification of the amorphous phase [4]. Still, the obtained by annealing a shock compacted amorphous Ni-49.12at %Ti powder. The increase of Ms seems to be Corresponding author. Fax: +43-1-4277-51316 unclear; it was proposed that enhanced nucleation is E-Jmail address: waitz(@ap univie. acat (T. Waitz) facilitated by internal stresses caused by the nanograins On leave from the Institute of advanced Materials. Ua state or by effects of the shock compression [7 Aviation Technical Univ Since contradictory results of the martensitic trans- Russian Federation 12 K Marksa Street, 450000 Ufa. formation presented in the literature may be attributed to 1359-6454/S30.00@ 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved 036

Martensitic phase transformations in nanocrystalline NiTi studied by TEM T. Waitz *, V. Kazykhanov 1 , H.P. Karnthaler Institute of Materials Physics, University of Vienna, Boltzmanngasse 5, A-1090 Vienna, Austria Received 16 June 2003; accepted 22 August 2003 Abstract By high pressure torsion (HPT) deformation almost complete amorphization is obtained in bulk Ni–50.3at.%Ti containing B190 martensite. During low temperature annealing tiny crystallites retained after the HPT deformation are acting as nuclei and trigger the nanocrystallization of B2 austenite. It is shown that the density of the nuclei is a function of the HPT strain and determines together with the annealing temperature the grain size of the nanocrystals ranging from 5 to 350 nm. Upon cooling the nano￾structures transform to B190 partially since the grain boundaries hinder the autocatalytic formation of martensite. The large transformation strains of B190 are reduced by very fine (0 0 1) compound twins. With decreasing grain size an increasing energy barrier arises and the martensitic transformation is completely suppressed in grains smaller than 60 nm. The R-phase transformation causing only small transformation strains is observed in grains between 15 and 60 nm. Whereas in grains below 15 nm B2 remains indicating that no transformation occurs at all.
2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: High pressure torsion; Bulk amorphous materials; Nanocrystallization; Martensitic phase transformation; High-resolution electron microscopy 1. Introduction In the intermetallic compound NiTi the shape mem￾ory effect and superelastic properties are related to a martensitic phase transformation being of considerable interest both from a scientific and a technological point of view. In addition, in NiTi amorphization can be in￾duced applying various solid state processes such as particle irradiation [1] and strong mechanical deforma￾tion by cold rolling [2] and mechanical alloying [3]. Recently severe plastic deformation by high pressure torsion (HPT) methods was applied to NiTi achieving amorphization; followed by a suitable heat treatment a nanocrystalline phase can obtained in bulk HPT alloys by devitrification of the amorphous phase [4]. Still, the cause of the nanocrystallization in the HPT deformed alloys remained unclear. As the martensitic transformation occurring in nanocrystalline NiTi is concerned there are contradic￾tory results presented in the literature. In Ni–49.8at.%Ti subjected to cold rolling crystal refinement and partial amorphization was achieved and the martensitic trans￾formation was suppressed even after cooling to )150 C [5,6]. The reason for the observed change in the trans￾formation seems to be unclear; it was proposed that dislocations induced by the cold rolling may stabilize the parent B2 austenite. Contrary to this, it was reported that the martensite start temperature Ms increases with decreasing grain size in a bulk nanocrystalline material obtained by annealing a shock compacted amorphous Ni–49.12at.%Ti powder. The increase of Ms seems to be unclear; it was proposed that enhanced nucleation is facilitated by internal stresses caused by the nanograins or by effects of the shock compression [7]. Since contradictory results of the martensitic trans￾formation presented in the literature may be attributed to Acta Materialia 52 (2004) 137–147 www.actamat-journals.com * Corresponding author. Fax: +43-1-4277-51316. E-mail address: waitz@ap.univie.ac.at (T. Waitz). URL: http://www.univie.ac.at/Materialphysik/EM. 1 On leave from the Institute of Advanced Materials, Ufa State Aviation Technical University, 12 K. Marksa Street, 450000 Ufa, Russian Federation. 1359-6454/$30.00
2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2003.08.036

T. Waitz et al. Acta Materialia 52(2004)137-14 stresses caused by lattice defects rather than to the ultra- of 40C/min in flowing nitrogen gas using a Perkin fine grain size it is the aim of the present study to carry out Elmer DSC 7 NiTialloy was investigated containing little lattice strains twin-jet polishing. The thinning was dome EM foils by a systematic investigation. Therefore a nanocrystalline The specimens were used to prepare Tenupol 3 and almost no dislocations and having a broad range of with a solution of 75% CH3oh and 25%HNO3(-22oC grain sizes(from about 5 to 350 nm). In a first step, the and 15 V). The different phases were analyzed by se- NiTi alloy showing the monoclinic B19 martensite phase lected area(SA) diffraction applying different beam di at room temperature(rT) was subjected to HPT defor- rections(BD)in a TEM(Philips CM 200 operating at mation achieving a bulk amorphous alloy. As compared 200 kV). Afterwards a hRTEM analysis was carried out to mechanically amorphized powders no further com- using a Philips CM30 ST(operating at 250 and 300 kv) paction or densification is necessary and contrary to cold equipped with a Gatan slow scan CCD camera rolling almost complete amorphization can be achieved by the HPT method. In a second step, a nanocrystalline microstructure was obtained by annealing the alloy close 3. Experimental results to the crystallization temperature. To investigate the nanocrystallization, two different degrees of HPT defor- 3. 1. The nanostructured amorphous phase after HPT of mation were carried out and the microstructures were different strains analysed carefully prior and after the devitrification. It ddition at rt the transformation induced microstruc Fig. I shows a TEM study of a specimen immediately ture and its dependence on the grain size was investigated after the hPt deformation at a strain of s= 6.7.As seen in detail using both transmission electron microscopy in Fig. 1(a)(bright field image) the specimen contains (TEM) and high resolution transmission electron heterogeneously distributed nanocrystals(spots of dark microscopy(HRTEM) methods contrast) that are embedded in an amorphous matrix and have a size between 5 and 30 nm. As analyzed by HRTEM and sa diffraction methods most of the larger 2. Experimental procedure nanocrystals contain B19 martensite whereas the smaller crystallites have the B2 structure. The contrast is reversed In the present study a binary NiTi alloy with a in the dark held image of Fig. I(b)that was achieved by nominal composition Ni-50.at. %Ti was used. The de- placing an objective aperture over part of the( 110)B2, tails of the alloy preparation and the transformation(11)Big and (020)Big diffraction rings of the nano- temperatures are given by [8]: the initial coarse grained crystals. a band shaped area that contains only a very alloy shows a single step transformation from B2 aus- small volume fraction(<1%)of crystallites that have a tenite to monoclinic B19 martensite(the martensite diameter in the range of 5-15 nm is marked by L; an area finish temperature Mr is in the range from 22 to 45C containing a higher density of mainly larger crystals is depending on the thermo-mechanical treatment prior to marked by H. In the diffraction pattern of Fig. I(c)the The alloy was quenched from 800oC in water below by the rather sharp diffraction rings of the crystallo o he transformation) amorphous phase gives rise to diffuse rings superimpose r and used to prepare HPT discs applying 10 turns at a seen g. 2 almost the entire volume of the pressure of 6 GPa. The as-processed HPT discs had a specimens deformed at the higher strain S=7.3 is diameter of 12 mm and a thickness of about 0. 2 mm. amorphous. A low density (less than about 1%)of From the hPt discs specimens with a diameter of 2.3 nanocrystals is observed in bright and dark field images mm were punched by spark erosion using very low (cf. Figs. 2(a) and(b)). Their analysis shows that the power to avoid any heating. To study the influence of nanocrystals are heterogeneously distributed; areas are strain on the amorphization specimens were taken at observed containing a higher density of nanocrystal two different distances from the centre of the hPt discs:(e.g. near h) whereas other areas are almost free of them about 2.5 and 4.3 mm; this corresponds to true loga-(e.g. near L). The crystallites have a size of less than rithmic strains S of about 6.7 and 7.3, respectively, at the about 15 nm and using HRTEM methods some of them ntral area of the TEM specimens. To achieve different with a diameter of only 3 nm were analyzed. In the grain sizes two different heat treatments were carried diffraction pattern(see Fig. 2(c))diffuse rings are ob- out: the specimens were annealed either at 340C for 5 h served; the radius of the inner, bright ring correspond or at 450C for I h under vacuum. The heat treatment to 4.7 nm and therefore to the reflection gllo of the was followed by cooling to RT and by quenching either NiTi B2 lattice; additional rings being weak and broad into methanol at a temperature of -25C or into liqu correspond to g211 and g220(as indicated in Fig. 2(c)) nitrogen. The peak temperature Tp and the onset tem- The analysis of several diffraction patterns shows that perature Tx of the crystallization were measured by frequently the first diffuse ring is superimposed by weak differential scanning calorimetry(DSC)at a heating rate(1 10)B2 reflection spots of the nanocrystals In addition

stresses caused by lattice defects rather than to the ultra- fine grain size it is the aim of the present study to carry out a systematic investigation. Therefore a nanocrystalline NiTi alloy was investigated containing little lattice strains and almost no dislocations and having a broad range of grain sizes (from about 5 to 350 nm). In a first step, the NiTi alloy showing the monoclinic B190 martensite phase at room temperature (RT) was subjected to HPT defor￾mation achieving a bulk amorphous alloy. As compared to mechanically amorphized powders no further com￾paction or densification is necessary and contrary to cold rolling almost complete amorphization can be achieved by the HPT method. In a second step, a nanocrystalline microstructure was obtained by annealing the alloy close to the crystallization temperature. To investigate the nanocrystallization, two different degrees of HPT defor￾mation were carried out and the microstructures were analysed carefully prior and after the devitrification. In addition, at RT the transformation induced microstruc￾ture and its dependence on the grain size was investigated in detail using both transmission electron microscopy (TEM) and high resolution transmission electron microscopy (HRTEM) methods. 2. Experimental procedure In the present study a binary NiTi alloy with a nominal composition Ni–50.3at.%Ti was used. The de￾tails of the alloy preparation and the transformation temperatures are given by [8]: the initial coarse grained alloy shows a single step transformation from B2 aus￾tenite to monoclinic B190 martensite (the martensite finish temperature Mf is in the range from 22 to 45 C depending on the thermo-mechanical treatment prior to the transformation). The alloy was quenched from 800 C in water below Mf and used to prepare HPT discs applying 10 turns at a pressure of 6 GPa. The as-processed HPT discs had a diameter of 12 mm and a thickness of about 0.2 mm. From the HPT discs specimens with a diameter of 2.3 mm were punched by spark erosion using very low power to avoid any heating. To study the influence of strain on the amorphization specimens were taken at two different distances from the centre of the HPT discs: about 2.5 and 4.3 mm; this corresponds to true loga￾rithmic strains S of about 6.7 and 7.3, respectively, at the central area of the TEM specimens. To achieve different grain sizes two different heat treatments were carried out: the specimens were annealed either at 340 C for 5 h or at 450 C for 1 h under vacuum. The heat treatment was followed by cooling to RT and by quenching either into methanol at a temperature of )25 C or into liquid nitrogen. The peak temperature Tp and the onset tem￾perature Tx of the crystallization were measured by differential scanning calorimetry (DSC) at a heating rate of 40 C/min in flowing nitrogen gas using a Perkin– Elmer DSC 7. The specimens were used to prepare TEM foils by twin-jet polishing. The thinning was done in a Tenupol 3 with a solution of 75% CH3OH and 25% HNO3 ()22 C and 15 V). The different phases were analyzed by se￾lected area (SA) diffraction applying different beam di￾rections (BD) in a TEM (Philips CM 200 operating at 200 kV). Afterwards a HRTEM analysis was carried out using a Philips CM30 ST (operating at 250 and 300 kV) equipped with a Gatan slow scan CCD camera. 3. Experimental results 3.1. The nanostructured amorphous phase after HPT of different strains Fig. 1 shows a TEM study of a specimen immediately after the HPT deformation at a strain of S ¼ 6:7. As seen in Fig. 1(a) (bright field image) the specimen contains heterogeneously distributed nanocrystals (spots of dark contrast) that are embedded in an amorphous matrix and have a size between 5 and 30 nm. As analyzed by HRTEM and SA diffraction methods most of the larger nanocrystals contain B190 martensite whereas the smaller crystallites have the B2 structure. The contrast is reversed in the dark field image of Fig. 1(b) that was achieved by placing an objective aperture over part of the h110iB2, h
111iB190 and h020iB190 diffraction rings of the nano￾crystals. A band shaped area that contains only a very small volume fraction (<1%) of crystallites that have a diameter in the range of 5–15 nm is marked by L; an area containing a higher density of mainly larger crystals is marked by H. In the diffraction pattern of Fig. 1(c) the amorphous phase gives rise to diffuse rings superimposed by the rather sharp diffraction rings of the crystallites. As seen in Fig. 2 almost the entire volume of the specimens deformed at the higher strain S ¼ 7:3 is amorphous. A low density (less than about 1%) of nanocrystals is observed in bright and dark field images (cf. Figs. 2(a) and (b)). Their analysis shows that the nanocrystals are heterogeneously distributed; areas are observed containing a higher density of nanocrystals (e.g. near H) whereas other areas are almost free of them (e.g. near L). The crystallites have a size of less than about 15 nm and using HRTEM methods some of them with a diameter of only 3 nm were analyzed. In the diffraction pattern (see Fig. 2(c)) diffuse rings are ob￾served; the radius of the inner, bright ring corresponds to 4.7 nm
1 and therefore to the reflection g110 of the NiTi B2 lattice; additional rings being weak and broad correspond to g211 and g220 (as indicated in Fig. 2(c)). The analysis of several diffraction patterns shows that frequently the first diffuse ring is superimposed by weak h110iB2 reflection spots of the nanocrystals. In addition 138 T. Waitz et al. / Acta Materialia 52 (2004) 137–147

T. Waitz et al. Acta Materialia 52(2004)137-147 200nm 200nm L H ⊙ 5nm g Fig. 1. Ni-50.3at %Ti after HPT deformation; strain S=6.7.(a) TEM 0.3at %Ti after HPt deformation; strain S= 7.3.(a) TEM bright field image. The amorphous matrix contains retained nano- ous matrix contains a volume fraction als showing dark contrast. P marks a coarse particle of the Ti2Ni of retained nanocrystals of less than about I %.(b) Crystallites having a lattice.(b) TEM dark field image of the heterogeneously distributed diameter of less than about 15 nm she ystallites showing bright contrast(H and L mark areas containing a corresponding TEM dark field image. H marks a higher density of high and low density of them).(c) In the diffraction pattern broad nanocrystals whereas in the area marked L there are almost no fuse rings of the amorphous phase are superimposed with rings nanocrystals. (c)Sa diffraction pattern showing broad diffuse rings of containing B2 and B19 diffraction spots of the nanocrystallites. he amorphous phase having radii that correspond to the length of the diffraction vectors(1 10),(211>and(220) of the B2 lattice

Fig. 2. Ni–50.3at.%Ti after HPT deformation; strain S ¼ 7:3. (a) TEM bright field image. The amorphous matrix contains a volume fraction of retained nanocrystals of less than about 1%. (b) Crystallites having a diameter of less than about 15 nm show up as bright spots in the corresponding TEM dark field image. H marks a higher density of nanocrystals whereas in the area marked L there are almost no nanocrystals. (c) SA diffraction pattern showing broad diffuse rings of the amorphous phase having radii that correspond to the length of the diffraction vectors h110i, h211i and h220i of the B2 lattice. Fig. 1. Ni–50.3at.%Ti after HPT deformation; strain S ¼ 6:7. (a) TEM bright field image. The amorphous matrix contains retained nano￾crystals showing dark contrast. P marks a coarse particle of the Ti2Ni lattice. (b) TEM dark field image of the heterogeneously distributed crystallites showing bright contrast (H and L mark areas containing a high and low density of them). (c) In the diffraction pattern broad diffuse rings of the amorphous phase are superimposed with rings containing B2 and B190 diffraction spots of the nanocrystallites. T. Waitz et al. / Acta Materialia 52 (2004) 137–147 139

T. Waitz et al. Acta Materialia 52(2004)137-14 to the amorphous rings weak(200)B2 reflections were observed. It should be mentioned that a volume fraction of about 5% corresponds to coarse spherical particles of the Ti2 Ni lattice that have survived the hpt in both cases S=6.7 and 7.3(cf. P in Fig. 1(a) 3. 2. Nanocrystallization after different heat treatments The onset of the crystallization was measured by DSC (cf. Fig 3). The analysis of the dSc curve of the HPT specimen with the lower strain(S= 6.7) yields that both the crystallization temperature and the crystall zation enthalpy are lower than in the case of s=7.3. Tx (Tp)are 352C(374C)and 362C(379C)in the case of S=6.7 and 7.3, respectively. The crystallization en- harpies△H67and△H73 are about-1.4and-1.7kJ mol corresponding to S=6.7 and 7.3, respectively Fig. 4 shows TEM bright field images of annealed specimens. Figs. 4(a) and(b)correspond to the defo mations S=6.7 ands=7.3, respectively; in addition the R pecimens were isothermally annealed at a temperature T=352°c 0,0 S=6.7 -0.5 374°C 00 Temperature['CI 0.5 T=362°c 0,0 0.5 Fig 4 Nanocrystalline phase formed after isothermal annealing of HPT deformed amorphous Ni-503at %Ti. TEM bright field images. T=379°c (a)S=6.7 after annealing at 340C for 5 h. Most of the grains are -1.0 smaller than about 50 nm containing both B2 and R-phase.(b 7.3 after annealing at 340C for 5 h. Frequently areas are Temperature[C] observed that contain mainly smaller grains (e.g. near S) and are adjacent to areas where larger gra dominating (e. g. near M Fig. 3. DSC curves of the crystallization of HPT deformed Ni- B2 phase and R-phase are found in the smaller grains near S. R 503at%Ti (heating rate 40C/min). (a)S=6.7. An exothermic peak marks a grain of the R-phase. Martensite occurs in the grains near occurs at Tp=374C; the crystallization temperature T=352C is M having a diameter of about 120 nm.(c)S=7.3 after annealing at indicated.(b)S=7.3. The crystallization occurs at a higher tempera- 450.C for I h. Almost all grains larger than about 150 nm contain ture(p=379°C,Tx=362°C

to the amorphous rings weak h200iB2 reflections were observed. It should be mentioned that a volume fraction of about 5% corresponds to coarse spherical particles of the Ti2Ni lattice that have survived the HPT in both cases S ¼ 6:7 and 7.3 (cf. P in Fig. 1(a)). 3.2. Nanocrystallization after different heat treatments The onset of the crystallization was measured by DSC (cf. Fig. 3). The analysis of the DSC curve of the HPT specimen with the lower strain (S ¼ 6:7) yields that both the crystallization temperature and the crystalli￾zation enthalpy are lower than in the case of S ¼ 7:3. Tx (Tp) are 352 C (374 C) and 362 C (379 C) in the case of S ¼ 6:7 and 7.3, respectively. The crystallization en￾thalpies DH6:7 and DH7:3 are about )1.4 and )1.7 kJ/ mol corresponding to S ¼ 6:7 and 7.3, respectively. Fig. 4 shows TEM bright field images of annealed specimens. Figs. 4(a) and (b) correspond to the defor￾mations S ¼ 6:7 and S ¼ 7:3, respectively; in addition the specimens were isothermally annealed at a temperature Fig. 3. DSC curves of the crystallization of HPT deformed Ni– 50.3at.%Ti. (heating rate 40 C/min). (a) S ¼ 6:7. An exothermic peak occurs at Tp ¼ 374 C; the crystallization temperature Tx ¼ 352 C is indicated. (b) S ¼ 7:3. The crystallization occurs at a higher tempera￾ture (Tp ¼ 379 C, Tx ¼ 362 C). Fig. 4. Nanocrystalline phase formed after isothermal annealing of HPT deformed amorphous Ni–50.3at.%Ti. TEM bright field images. (a) S ¼ 6:7 after annealing at 340 C for 5 h. Most of the grains are smaller than about 50 nm containing both B2 and R-phase. (b) S ¼ 7:3 after annealing at 340 C for 5 h. Frequently areas are observed that contain mainly smaller grains (e.g. near S) and are adjacent to areas where larger grains are dominating (e.g. near M). B2 phase and R-phase are found in the smaller grains near S. R marks a grain of the R-phase. Martensite occurs in the grains near M having a diameter of about 120 nm. (c) S ¼ 7:3 after annealing at 450 C for 1 h. Almost all grains larger than about 150 nm contain martensite. 140 T. Waitz et al. / Acta Materialia 52 (2004) 137–147

T. Waitz et al. Acta Materialia 52(2004)137-147 of 340C for 5 h followed by cooling to RT and quenching to -25C. Fig. 4(c) shows a specimen with S=7.3 annealed at a higher temperature(450C)for 1 h. In all cases the crystallization is complete since the diffraction patterns do not contain diffuse rings any more corresponding to the amorphous phase. Sharp and rather flat grain boundaries were observed both by TEM mean 30nm bright field and HRTEM images. It is important to point out that within the grains almost no dislocations were observed and most of the grains show only weak strain contrast Additional tem bright and dark field taken to measure the size distribution of the grains; the results are summarized in Fig. 5. In the case of S= 6.7 020406080100120140 annealed at 340C the grains have a diameter in the Grain size [nm] range of about 5-90 nm; only few of them are larger than about 50 nm(as shown in Fig. 5(a). In the case of S=7.3 annealed at 340C there is a broad range of grain diameters ranging from 5 to 140 nm(cf Fig. 5(b) Their distribution seems to be non-uniform since fre zza small grains quently areas are observed mainly containing grains of mean 25nm small diameters(mean diameter of about 25 nm, e.g. near S in Fig 4(b))whereas other areas that are adjacent large grains to them contain mainly larger grains(mean diameter mean 70nm about 70 nm, e.g. near M in Fig. 4(b)). In the case of S=7.3 and an annealing temperature of 450C most of LL 10 the grains are larger than about 100 nm(cf Fig. 5(c) Sa diffraction patterns were taken to analyze the H tt crystalline phases occurring in the grains of different di ameter. It is interesting to note that grains smaller than 020406080100120140 about 15 nm show reflections of the B2 phase only. When Grain size [nm the grain size is between 15 and 60 nm reflections of the B2 phase and the r-phase were observed whereas no reflections corresponding to the martensite were en- countered. When the grains are larger than about 60 nm 25 they contain R-phase(cf. the grain marked by R in Fig 4(b)and b19 martensite(cf. the area near M in Fig 4(b)) and in this case the B2 phase is hardly observed d15 mean 160nm Finally, grains larger than about 150 nm contain mainly martensite(cf. Fig. 4(c)). The volume fraction trans- formed to martensite by quenching to -25C was esti- mated to be less than about 30%(cf fig 4(b)and more than about 80%(cf Fig 4(c))in the specimens having a maximum grain size of about 140 and 350 nm(cf Figs. r (b)and (c), respectively. It should be noted that no 050100150200250300350 martensite could be detected in grains smaller than about 60 nm even when the specimens were quenched in liquid (c) Grain size [nm] nitrogen. In this case the volume fraction of R-phase Fig. 5. Histograms of the size distributions of the grains after crystalli seems to increase and only very small grains (less than zation (cf Fig 4(a).)S=6.7 after annealing at 340 C for 5 h is leading about 15 nm diameter)contain residual austenite. The to a mean grain size of 30 nm (b)s=7.3 after annealing at 340C for 5 results of this analysis are summarized in Table I h. The dashed bars correspond to areas containing mainly smaller and 3.3. The martensitic transformations in the nanostructures and the b19 martensite were analyzed in detail using In the specimens annealed at 340C for 5 h followed both SA diffraction and HRTEM methods. As illus- by cooling to RT and quenching to -25C the R-phase trated in Fig. 6 the grains containing the R-phase and

of 340 C for 5 h followed by cooling to RT and quenching to )25 C. Fig. 4(c) shows a specimen with S ¼ 7:3 annealed at a higher temperature (450 C) for 1 h. In all cases the crystallization is complete since the diffraction patterns do not contain diffuse rings any more corresponding to the amorphous phase. Sharp and rather flat grain boundaries were observed both by TEM bright field and HRTEM images. It is important to point out that within the grains almost no dislocations were observed and most of the grains show only weak strain contrast. Additional TEM bright and dark field images were taken to measure the size distribution of the grains; the results are summarized in Fig. 5. In the case of S ¼ 6:7 annealed at 340 C the grains have a diameter in the range of about 5–90 nm; only few of them are larger than about 50 nm (as shown in Fig. 5(a)). In the case of S ¼ 7:3 annealed at 340 C there is a broad range of grain diameters ranging from 5 to 140 nm (cf. Fig. 5(b)). Their distribution seems to be non-uniform since fre￾quently areas are observed mainly containing grains of small diameters (mean diameter of about 25 nm, e.g. near S in Fig. 4(b)) whereas other areas that are adjacent to them contain mainly larger grains (mean diameter about 70 nm, e.g. near M in Fig. 4(b)). In the case of S ¼ 7:3 and an annealing temperature of 450 C most of the grains are larger than about 100 nm (cf. Fig. 5(c)). SA diffraction patterns were taken to analyze the crystalline phases occurring in the grains of different di￾ameter. It is interesting to note that grains smaller than about 15 nm show reflections of the B2 phase only. When the grain size is between 15 and 60 nm reflections of the B2 phase and the R-phase were observed whereas no reflections corresponding to the martensite were en￾countered. When the grains are larger than about 60 nm they contain R-phase (cf. the grain marked by R in Fig. 4(b)) and B190 martensite (cf. the area near M in Fig. 4(b)) and in this case the B2 phase is hardly observed. Finally, grains larger than about 150 nm contain mainly martensite (cf. Fig. 4(c)). The volume fraction trans￾formed to martensite by quenching to )25 C was esti￾mated to be less than about 30% (cf. fig 4(b)) and more than about 80% (cf. Fig. 4(c)) in the specimens having a maximum grain size of about 140 and 350 nm (cf. Figs. 5(b) and (c)), respectively. It should be noted that no martensite could be detected in grains smaller than about 60 nm even when the specimens were quenched in liquid nitrogen. In this case the volume fraction of R-phase seems to increase and only very small grains (less than about 15 nm diameter) contain residual austenite. The results of this analysis are summarized in Table 1. 3.3. The martensitic transformations in the nanostructures In the specimens annealed at 340 C for 5 h followed by cooling to RT and quenching to )25 C the R-phase and the B190 martensite were analyzed in detail using both SA diffraction and HRTEM methods. As illus￾trated in Fig. 6 the grains containing the R-phase and Fig. 5. Histograms of the size distributions of the grains after crystalli￾zation (cf. Fig. 4(a)).) S ¼ 6:7 after annealing at 340 C for 5 h is leading to a mean grain size of 30 nm. (b) S ¼ 7:3 after annealing at 340 C for 5 h. The dashed bars correspond to areas containing mainly smaller and larger grains (mean 25 and 70 nm), respectively. (c) S ¼ 7:3 after an￾nealing at 450 C for 1 h. Almost all grains are larger than 100 nm. T. Waitz et al. / Acta Materialia 52 (2004) 137–147 141

T. Waitz et al. Acta Materialia 52(2004)137-14 Grain size dependence of the observed phases Grain size(nm) R- phase b19 15-60 60-150 (150 The nd B19 is indicated by estimate of the volume fraction of B19 is given in parentheses. S having a size larger than about 50-70 nm are frequently twinned and show some strain contrast: the width of the twins is about 20-50 nm(cf Fig. 6(a). Fig. 6(b)shows a HRTEM micrograph of two twin related R-phase 50nm variants RTI and RT2; they are contained in a grain aving a diameter of about 70 nm. The(1100)R lattice planes of both variants are indicated and they are cor responding to the(10 1)B2 and(01 1)B2 planes of the initial B2 lattice(in the case of the R-phase a hexagonal unit cell is used; BD runs along a common 2243R =11lB2 direction). The planar twin boundary is edge on in the TEM projection and runs along(110)B2. The 110R1 rhombohedral angle aR of the R-phase was measured tilting a grain containing two twin related variants to a common BD=[101OR(cf. Fig. 6(c). In this case, the diffraction pattern contains both the (0001 and (1210)R reflections. The corresponding lattice spacing doool and djzto were measured and used to calculate the (1100 ratio of the hexagonal unit cell parameters cR and ap cR/aR=1.39±0.0land1.38±0.01 leading to ar 2nm 88.5°±0.6°and87.8°±0.6° in the case of rtl and RT2, respectively Fig. 7 shows typical difiraction patterns correspond ing to b19 martensite. The pattern of Fig. 7(a) contains diffraction spots of two martensite twins MTI and MT2 that are related by a rotation of I around [0OlB (BD≡[10Mm≡[10Mm2). Most of the grains smaller x0001R2 than about 100-120 nm seem to contain only a single variant of the twinned martensite. In larger grains two 001RT1 or more martensite variants twinned on(00 1)Big planes 1210 T2 were frequently observed (cf. Fig. 7(b)). It should be noted that the reflections [11 l]MT1.2 and [11 1]MT1.2 are strongly elongated along 00 lBig and streaks are run RT1 ning along the [00 lBig direction indicating a very small width of the twins. In addition the streaks contain a regular sequence of closely spaced intensity maxima with an average spacing of 0.3 nm indicating a Fig. 6. Ni-50.3at %Ti Nanocrystalline structure: R-phase. (a) Bright fairly constant spacing of the twin boundaries of about field image of a twinned grain. Twin boundaries are indicated by arrows: 3 nm. Fig. 8(a)shows a TEM bright field image that can two twin related variants. The twin boundary and the( lio be used to measure the width of the twins the average f the variants rtl and rt2 are indicated (B width is d1= 1.7 and d2= 2.0 nm for MTI and MT2 [2243]RT1.2)(c)Diffraction pattern showing two twin related R-phase twins, respectively; the average relative twin period variants(BD=[1010 RT1.2)used to measure the rhombohedral angle A=d1/(d1+d2)is 0.46(cf. Fig 8(b). Almost all mar tensite grains contain(001)BIg twins that have a width using HRTEM images; the minimum width was mea- of less than 10 nm As illustrated in Fig. 9 the width of sured to be 4(002)BIg atomic planes(0.9 nm)only(near very narrow twins can be determined on an atomic level A in Fig. 9)

having a size larger than about 50–70 nm are frequently twinned and show some strain contrast; the width of the twins is about 20–50 nm (cf. Fig. 6(a)). Fig. 6(b) shows a HRTEM micrograph of two twin related R-phase variants RT1 and RT2; they are contained in a grain having a diameter of about 70 nm. The ð1
100ÞR lattice planes of both variants are indicated and they are cor￾responding to the ð1 0
1ÞB2 and ð0
1 1ÞB2 planes of the initial B2 lattice (in the case of the R-phase a hexagonal unit cell is used; BD runs along a common ½2 2
4 3R ½111B2 direction). The planar twin boundary is edge on in the TEM projection and runs along ð1
1 0ÞB2. The rhombohedral angle aR of the R-phase was measured tilting a grain containing two twin related variants to a common BD ½101 0 R (cf. Fig. 6(c)). In this case, the diffraction pattern contains both the ð0001ÞR and ð1
210ÞR reflections. The corresponding lattice spacing d0001 and d1
210 were measured and used to calculate the ratio of the hexagonal unit cell parameters cR and aR: cR/aR ¼ 1.39  0.01 and 1.38  0.01 leading to aR ¼ 88:5  0.6 and 87.8  0.6 in the case of RT1 and RT2, respectively. Fig. 7 shows typical diffraction patterns correspond￾ing to B190 martensite. The pattern of Fig. 7(a) contains diffraction spots of two martensite twins MT1 and MT2 that are related by a rotation of p around ½001 B190 (BD ½
110MT1 ½1
1 0MT2). Most of the grains smaller than about 100–120 nm seem to contain only a single variant of the twinned martensite. In larger grains two or more martensite variants twinned on ð001ÞB190 planes were frequently observed (cf. Fig. 7(b)). It should be noted that the reflections ½111MT1;2 and ½
1
1 1MT1;2 are strongly elongated along ½001 B190 and streaks are run￾ning along the ½001 B190 direction indicating a very small width of the twins. In addition, the streaks contain a regular sequence of closely spaced intensity maxima with an average spacing of  0:3 nm
1 indicating a fairly constant spacing of the twin boundaries of about 3 nm. Fig. 8(a) shows a TEM bright field image that can be used to measure the width of the twins. The average width is d1 ¼ 1:7 and d2 ¼ 2:0 nm for MT1 and MT2 twins, respectively; the average relative twin period k ¼ d1=ðd1 þ d2Þ is 0.46 (cf. Fig. 8(b)). Almost all mar￾tensite grains contain ð001ÞB190 twins that have a width of less than 10 nm. As illustrated in Fig. 9 the width of very narrow twins can be determined on an atomic level using HRTEM images; the minimum width was mea￾sured to be 4 ð002ÞB190 atomic planes (0.9 nm) only (near A in Fig. 9). Table 1 Grain size dependence of the observed phases Grain size (nm) B2 R-phase B190 150 * (>0.8) The occurrence of B2, R-phase and B190 is indicated by *. An estimate of the volume fraction of B190 is given in parentheses. Fig. 6. Ni–50.3at.%Ti. Nanocrystalline structure; R-phase. (a) Bright field image of a twinned grain. Twin boundaries are indicated by arrows; near S strain contrast is visible. (b) HRTEM image of a grain containing two twin related variants. The twin boundary and the ð1
100ÞRT1;2 lat￾tice planes of the variants RT1 and RT2 are indicated (BD ½2 2
4 3RT1;2) (c) Diffraction pattern showing two twin related R-phase variants (BD ½1 0
1 0 RT1;2) used to measure the rhombohedral angle. 142 T. Waitz et al. / Acta Materialia 52 (2004) 137–147

T. Waitz et al. Acta Materialia 52(2004)137-147 MT2 001 MT1,2 30nm M2 08642 汤MT2 Fig. 7. Ni-50.3at %Ti. Nanocrystalline structure; martensite. (a)SA diffraction pattern of a grain containing twinned martensite; (001) pound twinning is indicated by twin reflections(marked by mtI and width [nm reflections along 001 作wy如mdhm t % Ti. Nanocrystalline structure; martensite(a) TEM wins(BD=[IOMTI=[ 10Mr)(b)Two martensite variants MI and pound twins. The twin lamellae show bright and dark contrast(b) M2 that are twinned on(00 1)occur within one grain(BD=[110 M1,2). Histogram of the observed twin widths 4. Discussion rolling of NiTi a martensite microstructure refined by dislocation accumulation and strain induced twinning 4.1. HPT induced amorphization was observed containing numerous intersecting bands of nanocrystalline and amorphous phase arising locally by The results of Figs. I and 2 show that by hpt de- a shear strain instability [5, 6, 9, 10]. As compared to cold formation Ni-50.3at %Ti transforms from B19 mar- rolling the HPT strain is much higher and dominated by tensite to a nanostructured amorphous phase. It a shear deformation [11] causing shear bands and strong proposed that a localized deformation process proceeds crystal refinement [12]. Therefore, the HPt deformation the transformation by leading to a nanocrystalline of NiTi facilitates both the crystal refinement and the structure. The latter is fragmented by bands of an continuous accumulation of amorphous shear bands amorphous phase(cf Fig. 1; S=6.7). Similar, after cold When s increases the crystalline volume fraction

4. Discussion 4.1. HPT induced amorphization The results of Figs. 1 and 2 show that by HPT de￾formation Ni–50.3at.%Ti transforms from B190 mar￾tensite to a nanostructured amorphous phase. It is proposed that a localized deformation process proceeds the transformation by leading to a nanocrystalline structure. The latter is fragmented by bands of an amorphous phase (cf. Fig. 1; S ¼ 6:7). Similar, after cold rolling of NiTi a martensite microstructure refined by dislocation accumulation and strain induced twinning was observed containing numerous intersecting bands of nanocrystalline and amorphous phase arising locally by a shear strain instability [5,6,9,10]. As compared to cold rolling the HPT strain is much higher and dominated by a shear deformation [11] causing shear bands and strong crystal refinement [12]. Therefore, the HPT deformation of NiTi facilitates both the crystal refinement and the continuous accumulation of amorphous shear bands. When S increases the crystalline volume fraction Fig. 7. Ni–50.3at.%Ti. Nanocrystalline structure; martensite. (a) SA diffraction pattern of a grain containing twinned martensite; (0 0 1) compound twinning is indicated by twin reflections (marked byMT1 and MT2) related by a rotation of p around ½001. Streaks and elongated reflections along ½001MT1;2 are caused by the very small width of the twins (BD ½
110MT1 ½1
1 0MT2) (b) Two martensite variants M1 and M2 that are twinned on (0 0 1) occur within one grain (BD ½
110M1;2). Fig. 8. Ni–50.3at.%Ti. Nanocrystalline structure; martensite. (a) TEM bright field image of a grain containing a high density of (0 0 1) com￾pound twins. The twin lamellae show bright and dark contrast. (b) Histogram of the observed twin widths. T. Waitz et al. / Acta Materialia 52 (2004) 137–147 143

T. Waitz et al. Acta Materialia 52(2004)137-14 01 MT1 101)MT2 nm. Fig 9. Ni-50.3at %Ti. Nanocrystalline structure; martensite HRTEM image showing(00 1) compound twins having a minimum width of 4 lattice planes(0.9 nm)only (near A). The(001)MT1.2 twin boundary plane and the(10 1).] planes of the twin related martensite variants MTl and MT2 are indicated by a dashed and a full line, respectively(BD=[O1OMT12) gradually decreases until only isolated nanocrystals are 340C nanocrystallization occurs since the retained left embedded heterogeneously in an amorphous matrix. crystallites can act as heterogeneous nucleation sites Finally, caused by the plastic deformation of the leading to a high nucleation rate although the rate of amorphous matrix the retained nanocrystals dissolve in growth is low since Ta/ Tm=0.39 is small(Im melting he amorphous phase until at S=7.3 only a few very temperature). In agreement with the present conclusion small nanocrystals survive. Since most of them have the it was proposed that the devitrification process should B2 structure it is concluded that the austenite is more proceed with the largest nucleation rate and slowest stable against amorphization than the b19 martensite. growth rate to obtain a nanocrystalline structure; during This is in agreement with previous results [13]. In the low temperature annealing retained crystallites embed present case it is concluded that the nanocrystalline ded in the amorphous phase can act as nuclei since they debris contains austenite that was retained during the do not dissolve and can exceed the critical size for het thermally induced B2 to B19 transformation [14] prior erogeneous nucleation [17] to HPT. Still their might be another explanation: the It is concluded that the nucleation rate depends on observed B2 nanocrystals could be formed by a stress the HPt strain and decreases with increasing S(as can induced reverse transformation B19 to B2 during HPt be seen by an increase of Tx and Tp: cf. Fig 3). This is That might be caused by a local increase of temperature explained as follows: in the alloy deformed up to S=7.3 exceeding the austenite start temperature(As a 110C in the nucleation rate is lower since a only a small volume the present case). Similar results of a deformation in- fraction <1% of tiny crystallites <15 nm is retained(cf duced reverse transformation were reported by [15] Fig. 2). Contrary to this, in the case of s=6.7 nucle ation is triggered by numerous retained crystallites with 4.2. Nanocrystallization a diameter up to about 30 nm having a high nucleation potency(cf. Fig. I; a retained crystalline volume fraction The results of Figs. 3 and 4 show that during an- Ver of about 18% is deduced using Ver =1-(4H6.7/ ealing in HPT induced amorphous Ni-50. 3at %T1 △H7:3) loys heterogeneous polymorphous crystallization of the In the amorphous matrix the growth of the nuclei B2 phase occurs about 150C lower as compared to thin ceases when the advancing interfaces of neighboring amorphous NiTi films processed by sputtering or melt crystallites impinge on each other leading to grain spinning(Tr510C[16). It is concluded that the low boundaries. As the density of the nuclei increases their thermal stability of the amorphous phase formed by mean separation and therefore the final grain size will HPT is caused by the retained nanocrystalline debris decrease. Since the density of the nuclei is lower in the already at an annealing temperature as low as te, case of S=7.3 annealing at Ta=340C is leading to a triggering crystallization(see Figs. I and 2). Therefor larger grain size as compared to the alloy deformed up

gradually decreases until only isolated nanocrystals are left embedded heterogeneously in an amorphous matrix. Finally, caused by the plastic deformation of the amorphous matrix the retained nanocrystals dissolve in the amorphous phase until at S ¼ 7:3 only a few very small nanocrystals survive. Since most of them have the B2 structure it is concluded that the austenite is more stable against amorphization than the B190 martensite. This is in agreement with previous results [13]. In the present case it is concluded that the nanocrystalline debris contains austenite that was retained during the thermally induced B2 to B190 transformation [14] prior to HPT. Still their might be another explanation: the observed B2 nanocrystals could be formed by a stress induced reverse transformation B190 to B2 during HPT. That might be caused by a local increase of temperature exceeding the austenite start temperature (As  110 C in the present case). Similar results of a deformation in￾duced reverse transformation were reported by [15]. 4.2. Nanocrystallization The results of Figs. 3 and 4 show that during an￾nealing in HPT induced amorphous Ni–50.3at.%Ti al￾loys heterogeneous polymorphous crystallization of the B2 phase occurs about 150 C lower as compared to thin amorphous NiTi films processed by sputtering or melt spinning (Tx510 C [16]). It is concluded that the low thermal stability of the amorphous phase formed by HPT is caused by the retained nanocrystalline debris triggering crystallization (see Figs. 1 and 2). Therefore, already at an annealing temperature as low as Ta ¼ 340 C nanocrystallization occurs since the retained crystallites can act as heterogeneous nucleation sites leading to a high nucleation rate although the rate of growth is low since Ta=Tm ¼ 0:39 is small (Tm melting temperature). In agreement with the present conclusion it was proposed that the devitrification process should proceed with the largest nucleation rate and slowest growth rate to obtain a nanocrystalline structure; during low temperature annealing retained crystallites embed￾ded in the amorphous phase can act as nuclei since they do not dissolve and can exceed the critical size for het￾erogeneous nucleation [17]. It is concluded that the nucleation rate depends on the HPT strain and decreases with increasing S (as can be seen by an increase of Tx and Tp; cf. Fig. 3). This is explained as follows: in the alloy deformed up to S ¼ 7:3 the nucleation rate is lower since a only a small volume fraction <1% of tiny crystallites <15 nm is retained (cf. Fig. 2). Contrary to this, in the case of S ¼ 6:7 nucle￾ation is triggered by numerous retained crystallites with a diameter up to about 30 nm having a high nucleation potency (cf. Fig. 1; a retained crystalline volume fraction Vcr of about 18% is deduced using Vcr ¼ 1
ðDH6:7= DH7:3Þ. In the amorphous matrix the growth of the nuclei ceases when the advancing interfaces of neighboring crystallites impinge on each other leading to grain boundaries. As the density of the nuclei increases their mean separation and therefore the final grain size will decrease. Since the density of the nuclei is lower in the case of S ¼ 7:3 annealing at Ta ¼ 340 C is leading to a larger grain size as compared to the alloy deformed up Fig. 9. Ni–50.3at.%Ti. Nanocrystalline structure; martensite. HRTEM image showing (0 0 1) compound twins having a minimum width of 4 lattice planes (0.9 nm) only (near A). The ð001ÞMT1;2 twin boundary plane and the ð1 0
1ÞMT1;2 planes of the twin related martensite variants MT1 and MT2 are indicated by a dashed and a full line, respectively (BD ½010MT1;2). 144 T. Waitz et al. / Acta Materialia 52 (2004) 137–147

T. Waitz et al. I Acta Materialia 52(2004)137-147 to S=6.7 and annealed at the same temperature (cf. It should be noted that a suppression of a martensitic Figs. 4 and 5). In addition, when the retained crystallites phase transformation was also observed in small particles are not distributed uniformly a nanocrystalline phase and it was proposed that the fraction of particles con- with a heterogeneous distribution of the grain sizes taining a lattice defect suitable for heterogeneous nucle- arises after crystallization (as shown at S and M in ation is decreasing with decreasing grain size [24] Fig 4(b). Finally, as compared to an annealing tem- However, these statistical arguments seem to fail in the perature of 340C the growth rate is expected to be case of FeNi nanoparticles(size from 10-200 nm) since higher at Ta=450C(Ta/Tm=0.46) leading to larger the particles easily transform above RT although both grain sizes(compare Figs. 4(b) and(c)) experimental observations and calculations indicate that heterogeneous nucleation sites are not contained within 4.3. Phase transformations occurring in the nanocrystal he particles; therefore it was proposed that heteroge- neous nucleation sites are provided by the surfaces of the nanoparticles [25]. Similar, in the present case it is con Based on the results of the TEM analysis summarized cluded that in the nanocrystalline NiTi alloy the mar- in Table I it is concluded that in the nanocrystalline Ni- tensitic transformation is not suppressed by a lack of 603at %Ti alloy the martensitic transformation is sup- nucleation sites as the density of heterogeneous nucle- pre grain size. This ation sites provided by the grain boundaries may even transformed volume fraction decreases with decreasing increase with increasing grain boundary area and de- grain size and the onset of the martensitic transforma creasing n size ion is shifted towards lower temperatures. Therefore, Ms drops below the transformation temperature of the 43. 2. Constraints arising by the grain boundaries R-phase(TR)in grains smaller than about 150 nm It is proposed that the grain boundaries impose ading to a two step transformation from the B2 aus- constraints on the growth of the martensite confinin tenite via the R-phase to the B19 martensite. Finally, in the transformed volume fraction in the nanocrystalline grains smaller than about 60 nm the martensitic phase structure. A martensite plate nucleated within a grain transformation is completely suppressed. will be stopped at the grain boundaries acting as ob- Similar results of suppressing Ms were reported in the stacles To propagate the transformation the plate has to case of a cold rolled Ni-498Ti alloy containing a exert stresses that are sufficient to stimulate nucleation nanocrystalline phase [5]. a decrease of Ms below TR is and growth of favourable martensite variants in the also observed in coarse grained NiTi alloys and could be adjacent grains [22]. However, little stresses are expected caused by a high density of dislocations induced during to occur ahead of a small plate bound within an nano- rolling [18] or coherent precipitates induced by aging grain that has a size of less than 100 nm [26]. In addi- [19]. It was proposed that in a nanocrystalline NiTi alloy tion, as deduced from TEM observations in NiTi a the martensitic transformation could be suppressed in restriction was proposed for the autocatalytic nucleation different ways: by the introduction of elastic strains, by of self accommodating plate groups [27]. This is based lattice defects or by crystal refinement [5, 6]. In the on a twin relation of all the martensite variants in the present investigation almost no dislocations and little group which is possible only within a single grain since elastic strains were encountered in the grains(cf Fig 4). martensite variants occurring in adjacent grains have Therefore it is concluded that the grain refinement leads not the required orientation relationship. Therefore, in to the suppression of the transformation. In this context the present case it is concluded that in a volume con aspects concerning nucleation sites, grain boundaries, taining many small grains the ability of spreading the twinning and the R-phase transformation are discussed transformation by autocatalytic nucleation decreases in the following paragraphs with increasing grain boundary area in agreement with the observed decrease of the martensite volume fraction 4.3.1. Nucleation sites in the nanograins Specific dislocation configurations as dislocation walls 4.3.3.(001) Compound twinning or dislocation pile-ups are generally considered as possi As outlined by [28] a martensitic phase transforma- le nucleation sites [20, 21]; in coarse grained NiTi alloys tion will follow a path of almost complete accommo- there is direct experimental evidence that martensite is dation of the transformation shape strains. In coarse nucleating at dislocations tangles in the matrix and near grained NiTi two different mechanisms occurring at grain boundary dislocations[22]. Since in the present case different length scales facilitate the strain accommoda almost no dislocations occur within the nanograins it is tion: Firstly, twinning by(011) type II or (1 1 1)type I oncluded that the grain boundaries act as the heteroge twins(at scale of 30-100 nm) leading to an invariant neous nucleation sites. This is in agreement with experi- plane strain; secondly, at a scale far exceeding 100 nm mental results indicating that in NiTi the grain boundaries self-accommodation of groups of different habit plane can favour the martensitic transformation[23] variants arises [27, 29]. In the present case of the

to S ¼ 6:7 and annealed at the same temperature (cf. Figs. 4 and 5). In addition, when the retained crystallites are not distributed uniformly a nanocrystalline phase with a heterogeneous distribution of the grain sizes arises after crystallization (as shown at S and M in Fig. 4(b)). Finally, as compared to an annealing tem￾perature of 340 C the growth rate is expected to be higher at Ta ¼ 450 C ðTa=Tm ¼ 0:46Þ leading to larger grain sizes (compare Figs. 4(b) and (c)). 4.3. Phase transformations occurring in the nanocrystal￾line grains Based on the results of the TEM analysis summarized in Table 1 it is concluded that in the nanocrystalline Ni– 50.3at.%Ti alloy the martensitic transformation is sup￾pressed with decreasing grain size. This means, the transformed volume fraction decreases with decreasing grain size and the onset of the martensitic transforma￾tion is shifted towards lower temperatures. Therefore, Ms drops below the transformation temperature of the R-phase (TR) in grains smaller than about 150 nm leading to a two step transformation from the B2 aus￾tenite via the R-phase to the B190 martensite. Finally, in grains smaller than about 60 nm the martensitic phase transformation is completely suppressed. Similar results of suppressing Ms were reported in the case of a cold rolled Ni–49.8Ti alloy containing a nanocrystalline phase [5]. A decrease of Ms below TR is also observed in coarse grained NiTi alloys and could be caused by a high density of dislocations induced during rolling [18] or coherent precipitates induced by aging [19]. It was proposed that in a nanocrystalline NiTi alloy the martensitic transformation could be suppressed in different ways: by the introduction of elastic strains, by lattice defects or by crystal refinement [5,6]. In the present investigation almost no dislocations and little elastic strains were encountered in the grains (cf. Fig. 4). Therefore it is concluded that the grain refinement leads to the suppression of the transformation. In this context aspects concerning nucleation sites, grain boundaries, twinning and the R-phase transformation are discussed in the following paragraphs. 4.3.1. Nucleation sites in the nanograins Specific dislocation configurations as dislocation walls or dislocation pile-ups are generally considered as possi￾ble nucleation sites [20,21]; in coarse grained NiTi alloys there is direct experimental evidence that martensite is nucleating at dislocations tangles in the matrix and near grain boundary dislocations [22]. Since in the present case almost no dislocations occur within the nanograins it is concluded that the grain boundaries act as the heteroge￾neous nucleation sites. This is in agreement with experi￾mental results indicating that in NiTi the grain boundaries can favour the martensitic transformation [23]. It should be noted that a suppression of a martensitic phase transformation was also observed in small particles and it was proposed that the fraction of particles con￾taining a lattice defect suitable for heterogeneous nucle￾ation is decreasing with decreasing grain size [24]. However, these statistical arguments seem to fail in the case of FeNi nanoparticles (size from 10–200 nm) since the particles easily transform above RT although both experimental observations and calculations indicate that heterogeneous nucleation sites are not contained within the particles; therefore it was proposed that heteroge￾neous nucleation sites are provided by the surfaces of the nanoparticles [25]. Similar, in the present case it is con￾cluded that in the nanocrystalline NiTi alloy the mar￾tensitic transformation is not suppressed by a lack of nucleation sites as the density of heterogeneous nucle￾ation sites provided by the grain boundaries may even increase with increasing grain boundary area and de￾creasing grain size. 4.3.2. Constraints arising by the grain boundaries It is proposed that the grain boundaries impose constraints on the growth of the martensite confining the transformed volume fraction in the nanocrystalline structure. A martensite plate nucleated within a grain will be stopped at the grain boundaries acting as ob￾stacles. To propagate the transformation the plate has to exert stresses that are sufficient to stimulate nucleation and growth of favourable martensite variants in the adjacent grains [22]. However, little stresses are expected to occur ahead of a small plate bound within an nano￾grain that has a size of less than 100 nm [26]. In addi￾tion, as deduced from TEM observations in NiTi a restriction was proposed for the autocatalytic nucleation of self accommodating plate groups [27]. This is based on a twin relation of all the martensite variants in the group which is possible only within a single grain since martensite variants occurring in adjacent grains have not the required orientation relationship. Therefore, in the present case it is concluded that in a volume con￾taining many small grains the ability of spreading the transformation by autocatalytic nucleation decreases with increasing grain boundary area in agreement with the observed decrease of the martensite volume fraction. 4.3.3. (0 0 1) Compound twinning As outlined by [28] a martensitic phase transforma￾tion will follow a path of almost complete accommo￾dation of the transformation shape strains. In coarse grained NiTi two different mechanisms occurring at different length scales facilitate the strain accommoda￾tion: Firstly, twinning by h011i type II or ð1 1
1Þ type I twins (at scale of 30–100 nm) leading to an invariant plane strain; secondly, at a scale far exceeding 100 nm self-accommodation of groups of different habit plane variants arises [27,29]. In the present case of the T. Waitz et al. / Acta Materialia 52 (2004) 137–147 145

T. Waitz et al. Acta Materialia 52(2004)137-14 nanocrystalline NiTi the mean grain size(30-160 nm, cf to grain boundaries. Also the grain boundaries might Figs. 4 and 5)represents a length scale of the same order even trigger the formation of twins [34] as that of the twinning period and smaller than that of inally, it was proposed recently that during the r the self-accommodating groups in a coarse grained al- phase to B19 transformation(00 1)compound twinning loy. Therefore, in the nanograins the constraints of the could give rise to an invariant plane strain under the grain boundaries suppress the formation of self-accom- condition of a critical value aR 89o r2s case of proposed that (00 1)compound twinning leads to a critical angle but the agreement is better than in the ca twinning [29, 30]. In addition, a decrease of the twinning period d1+d2 decreases the elastic energy Awe(strain 4.3.4. R-phase transformation energy per unit transformed volume)[31]. However, It is concluded that the constraints of the grain with decreasing d the twin boundary area At and boundaries should have little effect on the B2 to R-phase therefore the twin boundary energy A,7t of the trans- transformation since in this case the transformation formed volume increase. From the observed very small shape strain is significantly smaller (roughly 1% as twinning period it can be concluded that the specific compared to 10% in the case of the B2 to B19 trans- twin boundary energy it of the(00 1) compound twins formation). This is in agreement with the present results must be very small. It should be mentioned that in the since on cooling the r-phase precedes the martensite in present case (cf. Fig. 9) the thinnest lamellae of the the small grains. The critical diameter of nuclei of the r compound twins contain two martensite unit cells (i.e. phase was estimated to be about 4-8 nm by in situ TEM four(002) lattice planes yielding d=0.9 nm). Since the analysis [36]. Therefore, the R-phase transformation separations of the twin boundaries are comparable to expected to be completely suppressed in very small the interatomic distance the nanotwinned martensite grains as shown in Table 1. In grains having a size in the may be regarded as an adaptive martensite phase as range from 15 to 50 nm small strains could be present proposed by [28]. Finally, experimental results found in prior to the transformation that may help to accom the literature [29, 30, 32] indicate that a variety of modate a single variant of the R-phase elastically In the constraints as dislocations, grain boundaries and pre- grains larger than 50 nm, however, the shape strains cipitates that may inhibit self-accommodation lead to have to be accommodated by twins that have a small the same preferred mode of the transformation involv- width to decrease Awe(cf. Fig. 6(a). Therefore as ex- ing(001)compound twinning to decrease the strain pected, caused by the constraints of the grain bound energy aries the width of the twins(20-50 nm) is significantly The experimental results show that in grains smaller smaller in the nanograins as compared to that of about than 60nm(cf. Table 1)no thermally induced martensitic 300 nm measured in coarse grained NiTi alloys[371 phase transformation occurs. It is proposed that this is caused by the decrease of self-accommodation with de- creasing grain size since per unit of transformed volume 5. Conclusions an increasing chemical driving force-Ag(0)including elastic(Awe)and surface ing to amorphization. Nanocrystalline debris is re- (At,)energy. The present conclusions are in agreement tained in the amorphous phase; the size and the with calculations of Agne carried out for ZrO2 nanopar- number of the retained crystallites is decreasing with ticles containing twinned martensite [33]. Similar to the increasing HPT strain present case smaller particles are more stable against the 2. The amorphous phase shows low thermal stability. It martensitic transformation than larger ones is concluded that heterogeneous crystallization is trig- It should be mentioned that previously (00 1)com gered by the retained crystallites that have survived pound twins of the martensitic phase were considered as the hpt deformation deformation twins only since they do not agree with the 3. Isothermal annealing is leading to nanocrystalline phenomenological theory of the B2 to B19 transfor- structures with grain sizes in the range of 5-350 nm mation in NiTi as no invariant habit plane can occur depending on the annealing temperature and the den [29]. In the nanocrystalline alloy the presence of an in- sity of the heterogeneous nucleation sites present in variant habit plane is less important for the transfor the amorphous phase. This leads to the conclusion mation since the bl9 phase will mainly occur attached that the strain of the hpt deformation determines

nanocrystalline NiTi the mean grain size (30–160 nm, cf. Figs. 4 and 5) represents a length scale of the same order as that of the twinning period and smaller than that of the self-accommodating groups in a coarse grained al￾loy. Therefore, in the nanograins the constraints of the grain boundaries suppress the formation of self-accom￾modating variants. As a consequence within the nano￾grains a different path of the transformation occurs that is leading to the required decrease of the transformation strains by the formation of very fine (0 0 1) compound twins (cf. Figs. 7–9). This is outlined as follows: it was proposed that (0 0 1) compound twinning leads to a smaller strain energy as compared to h011i type II twinning [29,30]. In addition, a decrease of the twinning period d1 + d2 decreases the elastic energy Dwe (strain energy per unit transformed volume) [31]. However, with decreasing d the twin boundary area At and therefore the twin boundary energy Atct of the trans￾formed volume increase. From the observed very small twinning period it can be concluded that the specific twin boundary energy ct of the (0 0 1) compound twins must be very small. It should be mentioned that in the present case (cf. Fig. 9) the thinnest lamellae of the compound twins contain two martensite unit cells (i.e. four (0 0 2) lattice planes yielding d ¼ 0:9 nm). Since the separations of the twin boundaries are comparable to the interatomic distance the nanotwinned martensite may be regarded as an adaptive martensite phase as proposed by [28]. Finally, experimental results found in the literature [29,30,32] indicate that a variety of constraints as dislocations, grain boundaries and pre￾cipitates that may inhibit self-accommodation lead to the same preferred mode of the transformation involv￾ing (0 0 1) compound twinning to decrease the strain energy. The experimental results show that in grains smaller than 60nm (cf. Table 1) no thermally induced martensitic phase transformation occurs. It is proposed that this is caused by the decrease of self-accommodation with de￾creasing grain size since per unit of transformed volume an increasing chemical driving force)Dgc (0) including elastic (Dwe) and surface (Atct) energy. The present conclusions are in agreement with calculations of Dgnc carried out for ZrO2 nanopar￾ticles containing twinned martensite [33]. Similar to the present case smaller particles are more stable against the martensitic transformation than larger ones. It should be mentioned that previously (0 0 1) com￾pound twins of the martensitic phase were considered as deformation twins only since they do not agree with the phenomenological theory of the B2 to B190 transfor￾mation in NiTi as no invariant habit plane can occur [29]. In the nanocrystalline alloy the presence of an in￾variant habit plane is less important for the transfor￾mation since the B190 phase will mainly occur attached to grain boundaries. Also the grain boundaries might even trigger the formation of twins [34]. Finally, it was proposed recently that during the R￾phase to B190 transformation (0 0 1) compound twinning could give rise to an invariant plane strain under the condition of a critical value aR 6 86:2 [29]. In the present case the martensitic transformation is preceded by the formation of R-phase as required for (0 0 1) compound twinning acting as lattice invariant strain. As deduced from Fig. 6(c) the measured RT values of aR (88.5  0.6 and 87.8  0.6) are somewhat larger than the calculated critical angle but the agreement is better than in the case of annealed coarse grained NiTi where aR P89 [35]. 4.3.4. R-phase transformation It is concluded that the constraints of the grain boundaries should have little effect on the B2 to R-phase transformation since in this case the transformation shape strain is significantly smaller (roughly 1% as compared to 10% in the case of the B2 to B190 trans￾formation). This is in agreement with the present results since on cooling the R-phase precedes the martensite in the small grains. The critical diameter of nuclei of the R￾phase was estimated to be about 4–8 nm by in situ TEM analysis [36]. Therefore, the R-phase transformation is expected to be completely suppressed in very small grains as shown in Table 1. In grains having a size in the range from 15 to 50 nm small strains could be present prior to the transformation that may help to accom￾modate a single variant of the R-phase elastically. In the grains larger than 50 nm, however, the shape strains have to be accommodated by twins that have a small width to decrease Dwe (cf. Fig. 6(a)). Therefore as ex￾pected, caused by the constraints of the grain bound￾aries the width of the twins (20–50 nm) is significantly smaller in the nanograins as compared to that of about 300 nm measured in coarse grained NiTi alloys [37]. 5. Conclusions 1. A NiTi alloy was subjected to HPT deformation lead￾ing to amorphization. Nanocrystalline debris is re￾tained in the amorphous phase; the size and the number of the retained crystallites is decreasing with increasing HPT strain. 2. The amorphous phase shows low thermal stability. It is concluded that heterogeneous crystallization is trig￾gered by the retained crystallites that have survived the HPT deformation. 3. Isothermal annealing is leading to nanocrystalline structures with grain sizes in the range of 5–350 nm depending on the annealing temperature and the den￾sity of the heterogeneous nucleation sites present in the amorphous phase. This leads to the conclusion that the strain of the HPT deformation determines 146 T. Waitz et al. / Acta Materialia 52 (2004) 137–147