MICROSTRUCTURE AND TENSILE PROPERTIES OF TWOBINARY NiTi-ALLOYS

E. Hornbogena, V. Mertingerb and D. WurzelaaRuhr-Universita¨t Bochum, Institut fu¨r Werkstoff, Lehrstuhl Werkstoffwissenschaft, Bochum,

GermanybUniversity of Miskolc, Institute of Material Science, Departement of Physical Metallurgy,Femtani Tanszk, Hungary

(Received March 22, 2000)(Accepted in revised form July 19, 2000)

1. Introduction

The temperatures of martensitic transformation are primarily controlled by thermodynamic equilibrium(T0, equ. 1), and therefore by the chemical composition of the transforming phase (austenite,b). Inaddition structural defects may modify the transformation behavior (temperature, hysteresis, nature ofphases), due to effects on nucleation and/or propagation of the transformation front. Defects include notonly dislocations (1d), antiphase domain boundaries (2d), but also dispersed small particles of a secondphase (3d), which all affect the course of the transformation. In addition defects (dislocations) canmodify the nature of the reaction by favoring the premartensitic phase, i.e. a two step reaction [1,2]. Thestress fields of dislocations may aid the formation of R-phase which implies smaller shear (gbR,0,01)than the final martensitic transformation (gba'0,2). The purpose of this work was to produce differentwell-defined microstructures of austeniteb, for an analysis of anomalous (transformation induced) andconventional tensile properties.

2. Materials and Experimental Methods

Two binary NiTi-alloys were chosen for this investigation. Both could be obtained as homogeneoussolid solution. The higher Ni-alloy was able to form Ni-rich precipitates (Ni4Ti3) by tempering theb-phase between 300 and 600°C, (Tables 1 and 2). The introduction of structural defects was achievedby the process of marforming (MF, TMF,Mƒ), which renders the alloys untransformable. An additionalaging treatment leads to reformation of defectb-phase (dislocations, subboundaries) and in addition toformation of particles in alloy B (Table 2).

Tensile tests were conducted with all nine microstructures (Table 2), to determine data on thepseudo-elastic plateau and conventional yield stress. The latter was obtained at 200°C.Aƒ, a temper-ature at which no stress-induced transformation occurs and aging effects are still negligible. Thetemperature dependence of the pseudo-yield stress was determined subsequently at much lower stressesto avoid true plastic deformation. TEM was required for the analysis of the essential microstructuralfeatures such as dislocations, subboundaries and especially small particles.

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3. Experimental Results

The DSC analysis shows the change of transformation temperatures and the occurence of a premarten-sitic reaction (Fig. 1). Figure 2a provides a survey on the effect of the nine different treatments on thetransformation temperaturesMm andAm.

Figure 2b provides data on particle sizes of the aged alloy B and the evidence for moderateprecipitation hardening by very small particles only.

Examples for typical forms of the stress strain curves are given in Figure 3, as well as the methodsof obtaining the data which are shown in Figure 4a. The ranges of reversibility of stress-inducedtransformation (pseudo-elastic temperature window) are also indicated 4a, b.

Typical LM-, and TEM-micrographics show the microstructures which are obtained by thermal andthermomechanical treatments (Figure 5, Table 2).

4. Discussion of the Results

The data obtained from the stress strain curves depend on a) chemical composition (or transformationtemperatures), b) precipitate structure (aging), c) defect structure (marforming1 aging), d) testtemperaturesT.Ms. (see Table 2)

Precipitation creates two effects simultanously. A decreasing Ni-content of the matrix,CNi andparticles acting as obstacles to propagation of the transformation front. The experimental results providedata on the lowering effect of Ni onMs: dMs/dCNi,0. Only the treatment which led to the largestparticle size shows the expected raise ofMs due to precipitation of Ni (Fig. 2, condition 6). All the otherheat treatments leave the peak transformation temperatures (Mm) rather uneffected, the course ishowever smeared out, (Fig. 1c). This can be explained by a lowering effect of the fine dispersoid onMs.An additional driving force (undercoolingDT) is required to move the reaction front. This can beanalysed by separating the effect of the chemical composition onT0, and of the precipitation hardeningof austeniteDsp, on the additional undercoolingDT:

Ms 5 T0 2 DT (1)

DMs 5dT0

dCNiDCNi 1

dDT

dDspDsp (2)

TABLE 1Chemical Compositions of Both Alloys and Phase Transformation Temperatures of the Recrystallized Initial

States

Chemical Composition, wt%

Alloy Ni Cu Fe C O Ti

A 55,15 ,0,03 ,0,05 0,076 0,05 RestB 55,6 ,0,03 ,0,05 0,055 0,08 Rest

Phase Transformation Temperatures, °C

Alloy M s Mm Mf As Am Af

A 21 13 3 34 47 56B 229 239 251 218 27 1

Ni-Ti ALLOYS172 Vol. 44, No. 1

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Ni-Ti ALLOYS 173Vol. 44, No. 1

whereDCNi andDsp are Ni-content, and hardening of austenite after a certain aging treatment (Fig. 2,Table 2) [3]. The effect of plastic deformation (MF) can be understood correspondingly. Only loweringof the transformation temperatures is expected and found, because the chemical composition staysconstant (DC 5 0), but work hardeningDsd takes place.

Figure 1. DSC—curves of a) condition 2, b) condition 9 and c) condition 4 (Table 2).

Figure 2. a) Transformation temperatures of the different conditions (Table 2). Conditions 1 and 2: show effect of chem.composition; conditions 2, 3, 4, 5, 6: precipitation effect with decreasing Ni-content and coherency; conditions 1 and 9: influenceof recovery; conditions 2, 7, 8: influence of recovery plus precipitation. b) Only the smallest precipitates increase the yieldstrength of the austenite (tested at 200 °C). Ni-rich precipitates lower the Ni-content of the matrix, therefore the slope ds/dTincreases as compared to the as betatized (Table 3). The smallest precipitates affect the slope in the opposite way, although theNi-content is changed in the same way as by the bigger precipitates.

Ni-Ti ALLOYS174 Vol. 44, No. 1

DMs 5dDT

dDsdDsd (3)

In the homogeneous alloy (A) work hardening is somewhat less effective in modifying the course of themartensitic transformation (and consequently in the slope of the peak) as compared to the alloy B withsimultaneous precipitation. The strain-induced dislocations are less impeded from rearranging byparticles as compared to alloy B.

There is clear evidence for a favorable effect of the mechanical treatment (introduction of latticedefects) on the creation of a premartensitic reaction (Fig. 1). A clear separation between the R- andM-reaction is always observed:

b 3 R 3 M, (4a)

while for the microstructures obtained simply by aging a calorimetric separation is mostly impossible(Figure 2)

b 3 ~R 1 M!. (4b)

Figure 3. a) Condition 2 and 7 loaded up to fracture at 25 and 200 °C. b) Condition 7 cycled at three different temperatures. At71 °C the stress to induce the transformation is just below the conventional yield stress. A residual strain is found after unloadingcaused by plastic deformation during the transformation.

Figure 4. a) Temperature dependence of PE yield stresssba, (example 7). b) Temperature window for PE of all 9 conditions (seeTable 2). If the residual strain is lower than 1% after unloading from 6%, the stress strain behavior is defined PE in this paper.Condition 7 and 9 have the same grain size, but in 7 precipitates were found. Therefore, in such microstructures precipitates leadto no further improvement.

Ni-Ti ALLOYS 175Vol. 44, No. 1

A broad temperature range indicates that the two reactions are overlapping.The stress required for the onset of the martensitic transformationsba can be described by a

modified Clausius-Clapeyron equation (Fig. 4):

dsba

dT5

Sba

eba

r 5hba

ebaTba

r (5)

Two sets of data support this relationship: the enthalpy of transformationhba and the transformationtemperatureTba follow from the DSC measurements, and the entropy of transformationsba isobtainable from the tensile test. These data are summarized in Table 3. A mass density (r) of 6.4 g/cm3

Figure 5. a) Optical micrograph of the as betatized state (alloy B). b) TEM-DF micrograph (122)Ni4Ti3 of the SHT and annealed(350 °C/1 h) alloy B with small precipitates Ni4Ti3. c) TEM-BF micrograph g5 (110) of the SHT and annealed (550 °C/100h) alloy B with larger precipitates Ni4Ti3 and dislocations. d) TEM-BF micrograph of the MF and annealed (550 °C/6 min) alloyB with small grains.

TABLE 3Calculated and Measured Entropy of Transformation (s), Respectively, Clausius Clapeyron Slopeds/dT

Obtained by Calorimetric (DSC) and Mechanical (M) Measurements

1 2 3 4 5 6 7 8 9

DSC M DSC M DSC M DSC M DSC M DSC M DSC M DSC M DSC M

Sba *1022

J/(gK)8,34 8,40 7,44 7,28 ? 6,05 8,75 8,41 8,13 7,74 7,47 7,25 ? 7,34 4,70 6,58 6,78 6,65

ds/dTMPa/K

7.9 8.0 7.1 6.9 ? 6.5 7.6 7.3 7.6 7.3 7.5 7.3 ? 7.2 5.0 7.0 6.8 6.7

Ni-Ti ALLOYS176 Vol. 44, No. 1

is used for the calculations. The difference between the measured mass densities of all conditions is lessthan the standard deviation of one condition.

In case of a two-stage reaction (equ. 4a) the entropy of transformation must be subdivided into twoterms:

hba

Tba

5 sba 5 sbR 1 sRa (6)

The tensile tests were not sensitive enough to separate two reactions, evidenced during cooling bycalorimetric methods (Fig. 3).

The analysis of the conventional yield stresssyb shows a moderate hardening effect due to smallprecipitates, while marforming creates considerable strengthening from 300 to 1000 MPa. This isimportant because the ratio of pseudo-elastic and true yield stress is regarded as the essential criterionfor stability of pseudo-elastic state under cyclic loading. The ratio should be small to avoid creation andsliding of dislocations during the stress induced transformation cycles:

0 ,sba

syb

, 1 (7)

This relationship is useful for the interpretation of the different temperature ranges in which pseudo-elasticity has been observed (Fig. 4). An increasing tendency is clearly shown:

betatized3 aged3marformed1 aged.

Favorable factors can be summarized as formation ofR-phase, work hardening of austenite, and to somedegree formation of very small coherent particles. Consequently it could be shown that the microstruc-ture created by thermo-mechanical treatment of a precipitation hardenable alloy is most favorable forcreation of a pseudo-elastic state. This structure evidently minimizes structural irreversibility due todefects which are created during martensitic cycles.

5. Conclusions

1. Besides chemical composition several microstructural aspect affect the transformation behavior ofNiTi-alloys.

2. Defects in austenite lower Ms and correspondingly raise the critical stress for the stress-inducedtransformation.

3. Dislocations may induce premartensitic R-phase and a wide transformation range, if it is overlappingwith martensitic transformation.

4. Small amounts of stress-induced residual martensite will favour thermal and stress-induced trans-formation.

5. Precipitation in austenite (b) creates an effect due to change in chemical composition as well as byparticles impeding propagation of martensite crystals. Only for overaged states (large particles) is theeffect of chemical composition dominating.

6. Microstructures created by thermo-mechanical treatments also affect (raise) the conventional yieldstress of austenitesb. A low ratio of the stress required for stress-induced transformationsba to sb

can be defined as a prerequisite for stable behavior of the material under thermo-mechanical fatigueconditions.

7. This behavior was obtained best by treatments consisting of marforming and tempering.

Ni-Ti ALLOYS 177Vol. 44, No. 1

Acknowledgment

We are grateful for support for one of the authors (V.M.) by the Union of German Academies ofScience. The work is part of SFB 459 (Sonderforschungsbereich der DFG, Formgeda¨chtnistechnik).

References

1. D. Treppmann, E. Hornbogen, and D. Wurzel, J. Phys. IV. C8(5), 596 (1995).2. D. Treppmann, E. Hornbogen, and D. Wurzel, Z. Metallkd. 89, 126 (1998).3. E. Hornbogen, Acta Metall. 33, 595 (1985).

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