Proceedings ofthe International Conference on Shape Memory and Superclastic Technologies
`May T-l 1, 20136. Pacific Grove. Caiiforoia, USA
`BrianBerg, M.R. Mitchell, and Jim Pmfi, Editors, 1) [43-154
`
`Copyright © 2003 ASM Intematicnalfi)
`All rights reserved,
`DOI: 10.13611cp1006mrstl43
`
`What Is The Big Deal About The Af Temperature?
`
`Ming H. Wu, Mark Polinsky and Neal Webb
`Memry Corporation, Bethei, Connecticut, U. S.A.
`
`Abstract
`
`The austenite transformation finish temperature, Af, has been specified as a key characteristic in
`product specifications for an increasing number of Nitinol medical devices. Many literatures also
`interpret the A; temperature as a sole parameter for predicting the Nitinol material properties. As
`a result, the industry is spending a significant amount of resources to test this transformation
`temperature as a means for product quality assurance.
`In the present study, cold drawn wires of
`Ti-55.8wt%Ni alloy were heat treated at a temperature of BSD-600°C for durations up to 120
`minutes. Mechanical properties were determined by tensile tests while the transformation
`temperatures were measured by differential scanning calorimetry (DSC) and bend and free
`recovery (BFR). The fatigue properties of selective specimens were also evaluated using a
`rotating beam fatigue test method. The purpose of the study is to evaluate the correlation
`between tensile mechanical properties and the Af temperatures measured using the various
`techniques. The results will also assess the influence of Af temperature on Nitinol fatigue life.
`
`Key Words
`
`Transformation temperatures, Af temperature, differential scanning calorimetry, bend and free
`recovery test, mechanical properties, rotating beam fatigue test
`
`Introduction
`
`Nitinol alloys when fully recrystallized after annealing at elevated temperatures greater than
`600°C exhibit a single stage martensitic transformation from the parent B2 to B19’ monoclinic
`martensite. However, for most medical devices and implants the Nitinol alloy is optimized for
`superelastic mechanical properties by cold working and subsequent heat—treatment at lower
`temperatures such that nano—sized subgrains, high density of dislocation and very fine Ni~rich
`precipitates are present
`in the material
`[1]. Transformation in this type of microstructure
`generally proceeds in two stages of BIZ a R—phase —> B19’ martensite.
`
`Defining the alloy formulation by specifying one of the transformation parameters measured by
`differential scanning calorimetry (DSC) of a fully annealed material is straight forward because
`hysteresis and transformation kinetics are relatively constant for a well defined microstructure
`[2]. However, defining the mechanical properties by specifying one of the transfon-nation
`parameters of a heat treated but not fully recrystallized alloy can be quite challenging because of
`the following issues:
`
`
`
`Edwards Exhibit 1022, p. 1
`
`

`

`Lack of linear correlation between mechanical properties and transformation parameters
`0
`Influence of the R-phase transformation on the mechanical properties
`0
`- Discrepancy among the transformation temperatures measured by different techniques
`
`In spite of the above difficulties, the Nitinol industry is use to adopting the Af temperature as a
`key parameter for specifying Nitinol devices [3]. The Food and Drug Administration (FDA) also
`endorsed this approach in its guidance for non—clinical tests for intravascular stents [4]. It is
`however recognized that materials having a constant Ag temperature may exhibit different
`mechanical and even different transformation characteristics [5].
`
`In a study on heat
`Several factors can influence the temperatures of transformation (TTR).
`treating an equiatomic NiTi alloy, Thoma et a1 [6] found that increasing reduction of cold work
`and lowering heat treat temperature always elevate the R—phase to austenite TTR but generally
`suppress the martensitic to austenite TTR. However, at low cold work reductions, the trend on
`the martensite to austenite TTR reverses to higher temperatures with lower heat
`treat
`temperatures, The martensitic transformation hysteresis is widened by increasing cold work and
`decreasing heat treat temperature, but the R—phase hysteresis is less sensitive to cold working or
`heat treat conditions. For Ni—rich compositions such as those commonly used for superelastic
`medical devices, precipitation of Ni—rich phases during heat treat elevates TTR as the matrix
`becomes depleted in Ni [7].
`
`To further clarify the validity of using the Af temperature for specifying Nitinol devices, the
`present study investigated the effects of heat treatment on mechanical properties and the Ar
`temperatures measured by both DSC and Bend and Free Recovery (BFR) methods. A group of
`specimens having a constant Af temperature was used for assessing their differences in plateau
`stresses, ultimate tensile stresses (UTS) and fatigue endurances.
`
`Experimental
`
`Cold drawn Ti-55.8wt%Ni wires of 0.020” in diameter were used for the present study. The as-
`drawn wires having an UTS of 257 Ksi were heat treated in a temperature range of 350—600°C
`for various durations up to 120 minutes. Mechanical properties were determined by tensile tests
`at 37°C while the TTR as defined in ASTM F2005-05 [8] were measured by differential scanning
`calorimetry (DSC) [9]. The functional Ar temperature was determined by a bend and free
`recovery (BFR) method of ASTM F2082—03 [10] in which procedure the test specimen was bent
`at -60°C before measuring the shape recovery on heating. The fatigue endurances of selective
`specimens with an as—drawn, heat treated surface finish and a BFR Af in the range of 27+l—2°C
`were studied using a rotating beam (RTB) fatigue test apparatus in a 37°C water bath.
`
`Results and Discussion
`
`DSC Transformation Temperatures
`DSC scans from 70°C to -100°C of specimens heat treated at 350-400"C did not detect any
`transformation with the only exception of R—phase peaks exhibited by the specimen heat treated
`at 400°C for 120 minutes, the TTR of which are Rsi49°C, RFZ6°C, As=33°C and Af=52°C.
`
`the DSC scans detected primarily the R—phase
`treated at 450°C,
`For specimens heat
`transformation. The TTR are plotted in Figure 1. Both R, and A; decrease while RI and A,
`increase rapidly with heat treat time signifying that the width of the transformation peak is
`becoming narrower and better defined with increasing heat treat duration as the cold work
`deformation is quickly annealed out. Although no martensitic peak was detected on cooling, the
`
`144
`
`
`
`[
`
`Edwards Exhibit 1022, p. 2
`
`

`

`heating peak starts to split after 30 minutes of heat treat suggesting some sluggish martensitic
`transformation during the cooling scan. There is a slight increase of TTR after 30 minutes of
`heat treat, which is probably induced by the precipitation of Ni—rich phases.
`
`As can be seen in Figure 2, the transformation of specimens heat treat at 500°C for less than 30
`minutes is very much similar to that of those heat treated at 450°C but with the R—phase
`transformation occurring at a lower temperature range, The transformation of those heat treated
`for extended durations longer than 30 minutes is characterized by split cooling and heating peaks,
`hence M5, Mf, R; and Rf’ can all be determined. Again, the TTRs increase with prolonging heat
`treat perhaps due to the increasing amount ofNi—rich precipitate. The hysteresis for both R—phase
`and martensite also decrease with time.
`
`DSC TTR of those heat treated at 550°C are plotted in Figure 3. The results are characterized by
`separated cooling peaks and a merged single—stage reverse transformation. The cold work strain
`appears to he annealed out very rapidly within the first few minutes as the RS and Rf quickly
`stabilize followed by rapid rises of all TTR due to the Ni—rich precipitation. The martensite
`hystersis narrows expeditiously as the material approaches the fully annealed state with time.
`
`1150C Heat Treat
`
`oRe QRffiRS’DAS I At
`
`Time (mine)
`
`Figure I .' DSC TTR ofspecimens heat treated at 450°C.
`
`
`
`
`
`
`
`
`
`Edwards Exhibit 1022, p. 3
`
`

`

`
`
`5090 Heat Treat
`
` —l—_l—_l—_l—_I—
`
`O
`
`20
`
`40
`
`ED
`60
`‘ Time(n1'ns)
`one {smells oMfARs‘ ARf' aAs lAf
`
`
`100
`
`120
`
`140
`
`Figure 2: DSC TTR ofspecimens heat treated at 500°C.
`
`BFR AfTemperature
`For specimens heat treated at 350°C for less than 30 minutes substantial superelasticity was
`observed during release fiom bending at -60°C. Hence no BFR A; could be determined. The Af
`temperatures measured by BFR of all other specimens are plotted in Figure 4. It is interesting to
`note that while the DSC scan failed to detect any transformation for specimens heat treated at
`350u400°C, the BFR method is able to detect shape recovery. During 350°C heat treat for longer
`than 30 minutes, the Af temperature decreases with time, suggesting that the annealing effect is
`dominant over the influence of precipitation as the precipitation kinetics is slow at 350°C. At
`400°C, the Af remains unchanged at about 20°C and gradually increases during the extended heat
`treat duration while at 450°C, the A; is relatively constant for up to 120 minutes. The constant Af
`over a period of heat treat duration may be caused by the balancing effect between annealing and
`the precipitation of Ni—rich intermetallics. The trends of BFR Af are generally consistent with
`those of the DSC Af.
`
`
`
`I46
`
`Edwards Exhibit 1022, p. 4
`
`

`

`
`
`550C Heat Treat
`
`
`
`Heat Treat Tlme (min)
`* 350C I 4000 £450C @5003 ' 550C
`
`Figure 4: The Af temperatures determined by BFR method ofSpecimens heat treated at 350-
`600°C.
`
`Correlation between DSC TTR and BFR Af
`The correlation between the DSC TTR and BFR Ar can be best demonstrated in Figure 5 which
`plots the DSC Rf’ and A; vs. the BFR At.
`It is clear that the DSC Af which measures the finish
`temperature for the transformation from R—phase to austenite is generally higher than the BFR Ar
`while the Rf’, the finish temperature for the transformation from martensite to R—phase, agrees
`very well with the BFR Ar.
`
`Tensile Properties of Specimens with BFR Ar of 27+l—2°C
`Tensile stress—strain curves of specimens having a BFR Af in the range of 27+l-2°C are shown in
`Figures 6-9 and the key parameters are summarized in Table. I. It is obvious from the data that
`
`Edwards Exhibit 1022, p. 5
`
`

`

`in spite of the similar BFR Af temperature of these specimens not only are the plateau stresses
`and the UTS vastly different, the mechanical characteristics among these specimens are also
`distinctively different.
`
`BFRAIVS DSCAf&Rf
`
`“"88
`
`DJCi!
`
`IX!C!
`
`EI—
`1']:
`06I].
`II:
`
`DmD
`
`iDSC—fif l DEC RTFHme
`
`so
`
`BFR Af (e)
`
`Figure 5.‘ A plot ofDSCAfande’ vs. the BFR Af
`
`Table I: Tensile rg‘rfl’ties 0 s ecimens havin a BFRA in the ran 2 o 27+/—20C.
`
`
`
`
`
`
`The specimen heat treated at 350°C for 60 minutes exhibits a work—hardened superelasticity with
`very high but slanted plateau, a smooth transition between linear elasticity and pseudoelasticity
`as well as a much smaller hysteresis when compared to those of other specimens with similar
`BFR Air. The Young’s modulus calculated based on the linearly elastic portion of the curve is 6.5
`Msi.
`
`Specimens heat treated at 400°C and above exhibit flat pseudoelastic plateau. The plateau
`stresses and the UTS decrease with increasing heat treat temperature. The pseudoelastic strain of
`the specimen heat treated at 400°C for 30 minutes (Figure 7) expands from 1.5% to 5.5%. The
`4% pseudoelastic strain is smaller than those of specimens heat treated at higher temperatures
`which typically expand from 1% to 6% (Figures 8 and 9). An R—phase plateau is present on the
`loading segments of specimens heat treated at 400°C and above.
`It is observed at 5-20 Ksi in
`Figure 7, at 20-40 Ksi in Figure 8 and becomes well defined at about 30 Ksi in Figure 9. The
`Young’s modulus calculated based on the linear elastic curve before the R—phase plateau in
`
`148
`
`Edwards Exhibit 1022, p. 6
`
`

`

`
`
`
`
`Figure 9 is about 11.2 Msi. Because the DSC Af temperature of 33°C is below the test
`temperature of 37°C, this value represents the true modulus for the B2 austenite.
`
`Most of the specimens heat treated at 450°C also have the BFR Af temperature in the range of
`27+/—2°C.
`The tensile stress-strain curves of specimens heat treated at 450°C for 2 and 60
`minutes are shown in Figures 10 and 11, respectively. The plateau stresses, UTS and BFR Af
`temperatures are compared with those of the one heat treated for 30 minutes (Figure 8) in Table
`2. Although these tensile curves are fundamentally similar and are characterized by the presence
`of R—phase plateau and flat martensite plateau, the upper and lower plateau stresses decrease
`substantially with prolonged heat treat time from 2 to 60 minutes. The UTS on the other hand is
`less influenced by the heat treat duration.
`
`350C160m1r: H eat Treat
`
`
`
`
`
`
`
`TenslleStress(Ksl)
`
`Tensile Strain (a)
`
`Figure 7: Tensile stress-strain curve oft: specimen heat treated at 400°Cjhr 30 minutes.
`
`149
`
`_|._LGUMDOD
`
`”'5
`Eto
`33..
`as
`
`2a5
`
`'—
`
`
`
`Edwards Exhibit 1022, p. 7
`
`

`

`
`
`
`
`
`
`4500i30min Heat “Eat
`
`g-H)
`E.
`ta
`
`Tensile Stra'n ('16:)
`
`Figure 9: Tensile stress—strain curve ofdispecimen heat treated at 500°Cfor 12.0 minutes.
`
`Correlation between TTR and Plateau Stresses
`In an attempt to understand the correlation between TTR and plateau stresses, specimens after
`heat treat at 500°C for 120 minutes were tensile tested at 10°C, 22°C, 37°C, 50°C and 70°C. The
`upper and lower plateau stresses are plotted in Figure 12 vs.
`the test temperature. The
`temperatures estimated by extrapolating the trend lines of Clausius-Clapeyron relationship to
`zero stress are —15°C and 33°C for UP and LP, respectively. According to the Clausius~
`Clapeyron relationship [11], these numbers should agree with M, and Rf’, respectively, as they
`are related to the martensite rather than the R—phase transfonnation. However, both are much
`higher than the athennal DSC M, of 63°C and Rf’ of 27°C in Figure 2. The discrepancies are
`most likely due to the reason that the present specimens are polycrystalline materials of fine grain
`structure instead of single crystals. Martensitie transferamtion in fine grain structure has a higher
`barrier of back stress to overcome [12].
`
`
`
`2 Bu
`
`:
`'I—
`
`Edwards Exhibit 1022, p. 8
`
`

`

`
`
`Rotating Beam Fatigue Endurance
`Fatigue endurances of specimens heat treated to have a BFR Ag of 27+/—2°C were tested using a
`rotating beam method to a maximum run-out cycle of 10,000,000. The results are tabulated in
`Table 3 and the S—N curves are plotted in Figure 13. No significant difference is obvious when
`comparing the low cycle fatigue life at strains equal or less than 1.0%. There is hOWever a
`noticeable distinction in high cycle fatigue life at 06-08% strain among the four groups of
`different heat treats. The group of the longest survival rate appears to be those heat treated at
`45 0°C for 30 minutes as all the specimens survived the 10M run—out cycle at both 0.8% and 0.6%
`strains. The group of 400°C heat treat for 30 minutes also performs well with only one specimen
`failed at 32,630 cycles at 0.8% strain. The worst perforning group is the one heat treated at
`500°C for 120 minutes having only one run-out at 0.6% strain.
`
`4508l2mit1 Heattreat
`
`EaI.
`(Ksl)
`
`
`
`
`
`::w
`E.N
`
`gE
`
`
`
`TensileSh'ess
`
`Tails Stain (it: )
`
`
`Figure 11: Tensile stress-strain curve ofa specimen heat treated at 450°Cfor 60 minutes.
`
`The superior fatigue endurance at O.6-0.8% strain of specimens heat treated at 400°C for 30
`minutes and those heat treated at 450°C for 30 minutes can be attributed to an R—phase plateau at
`a low stress level in combination with a high UTS after the heat treat. The R-phase plateau at
`
`
`
`Edwards Exhibit 1022, p. 9
`
`

`

`low stresses reduces the stress imposed on the specimen while a high UTS ensures that defects
`are less likely to be induced and accumulated during the cyclic deformation of rotating beam test.
`
`SDUCHZD min Heat Treat
` .8.
`
`0 s E T
`
`
`
`8 E(
`
`estTempem hire (C)
`
`I Upper Plateau .
`o Lever Plateau
`——Linear (Upper Plateau} —-Linear (Lower Plateau]
`
`
`Fignre 12: Upper and lower plateau stresses ofspecimens heat treated at 500°Cfor J20 minutes
`and tensile tested at various temperatures. The trend lines of linear regression represent the
`CIausius—Clapeyron relationship.
`
`Table 2: Tensile properties of 450°C heat treated specimens having a BFR AI in the range of
`
`27+/—2°C.
`HT Temp HT Time
`BFR A,
`UP
`1 LP
`UTS
`(09
`(min)
`1°C;
`(gsi)
`gigsi
`(Ksi
`48
`2
`28
`450
`80
`235
`36
`450
`72
`224
`30
`28
`29
`450
`60
`29
`66
`223
`
`4,448
`
`
`
`Table 3: Rotatin Beam Fati
`re Test Results
`Strain
`350°C
`400°C
`450°C
`500°C
`("/3)
`60mins
`30mins
`30mins
`120mins
`2 4
`1,586
`2,614
`3,265
`1,561
`1,519
`2,574
`2.4
`2,222
`2,285
`1,748
`2,968
`2.4
`2,367
`1,842
`5,911
`1.8
`3,620
`3,899
`
`1.8
`3,412
`3,707
`3,970
`2,309
`
`1.8
`3,459
`3,837
`3,980
`5,098
`
`1.0
`7,513
`16,051
`17,159
`16,594
`
`
`1.0
`16,847
`18,892
`19,162
`16,978
`
`18,70621,28213,9721.0 18,514
`
`
`
`
`
`
`Run-out
`1,038,990
`0.8
`32,630
`54,783
`
`Run-out
`0 .8
`5,048,640
`Run—cut
`39,482
`
`
`
`Run-out
`Run-out
`0.8
`Run-out
`41,668
`
`
`
`7,104,489
`Run-out
`0.6
`2,710,547
`Run-out
`
`
`Run—out
`7,926,772
`Run-out
`0.6
`Run—out
`
`
`Run-outRun-outRun-out0.6 Run-out
`
`
`
`
`152
`
`Edwards Exhibit 1022, p. 10
`
`

`

`
`
`RTB SN Curves
`
` BendStrain
`(“/a) 3
`
`
`
`M
`
`—I.
`
`05—
`.mrmno
`v
`i
`,
`..
`,
`.
`_
`I
`
`0 ‘i— I
`—l— J
`‘l—
`1.E+08
`‘l.E+03
`1.E+04
`1.E+05
`1.E+08
`1.E+07
` Cycle
`+ 3500I60m
`0
`3500160113 Run-Out —A— 400C130m
`A 4DOCI30m Run—Out ~+- 4500130m
`o 45OCI30m Run-Out
`—t— 50001120m
`o 5000I120m Run—Out
`
`
`
`Figure 13: Rotating beam S—N curves ofspecimens having a BFR Afof2 7+/-2 °C.
`
`Conclusions
`
`SJ”
`
`DSC method failed to detect any transformation of specimens heat treated at 350°C or those
`at 400°C for less than 120 minutes. The BFR method however was able to measure shape
`recovery with specimens heat treated at 350°C for longer than 30 minutes and those heat
`treated at 400°C and above.
`
`The R—phase TTR generally decrease with narrowing range (difference of RS-Rf) in the early
`stage before rising during extended heat treat duration. The R—phase TTR also decrease with
`increasing heat treat temperature.
`The martensitic TTR consistently increase with narrowing hysteresis with increasing heat
`treat temperature and duration.
`The BFR Ar temperature agrees very well with the DSC Rf’ instead of DSC Af temperature.
`For those having a BFR A; of 27+/-2°C, specimens heat treat at 350°C exhibit work-
`hardened pesudoelasticity while flat martensitic plateaus are present when heat treated at
`400°C and higher temperatures. The plateau stresses and UTS generally decrease with
`increasing heat treat temperature.
`Due to the fine grain structure, the temperatures determined by extrapolating UP and LP by
`Clausius-Clapeyron relationship to zero stress are much higher than the DSC Ms and Rf”,
`respectively.
`Among the specimens having a constant BFR Af of 27+l—2°C, the specimens heat treated at
`400-450°C performed much better than those heat treated at 350°C and 500°C in strain-
`controlled rotating beam fatigue tests due the presence of R—phase plateau at low stresses and
`high UTS.
`Materials having a similar BFR Af temperature can exhibit diversely different Young’s
`modulus, plateau streSSes, UTS as well as
`fatigue endurance.
`Therefore, BFR Af
`temperature alone can not be interpreted as the sole transformation parameter for specifying
`Nitinol devices.
`
`153
`
`
`
`Edwards Exhibit 1022, p. 11
`
`

`

`References
`
`[1] Tom, A., Van Geertruyden, W., Misiolek, W.Z., Han, H.D. and Wu, M.H., “Microstructural
`Characterization of Nitinol Stents”, Proceedings,
`International Conference on Shape
`Memory and Superelostic Technologies, Pacific Grove, California, 2006.
`[2] ASTM F2063-05, “Standard Specification for Wrought Nickel—Titanium Shape Memory
`Alloys for Medical Devices and Surgical Implants”, ASTM, 2005.
`[3] Pelton, A.R., DeCello, J. and Miyazaki, 8., “Optimization of Processing and Properties of
`Medical Grade Nitinol Wire”, Minimally Invasive Therapy and Allied Technologies, 9(1),
`2000, pp.107—118.
`[4] “Non—Clinical Tests and Recommended Labeling for Intravascular Stents and Associated
`Delivery Systems”, Guidance for Industgg and FDA Staff, US. Department of Health and
`Human Services, Food and Drug Administration, 2005, pp. 10.
`[5] Lopes, T.L., Gong, X-Y, and Trépanier, C., “Fatigue Performance of Nitinol Tubing with Af
`of 25°C”, Proceedings, International Conference on Shape Memory and Superelastic
`Technologies, Pacific Grove, California, 2003, pp. 311-320.
`[6] Thoma, P.E., Kao, M—Y., Fariabi, S. and AbuJudom, D.N., “The Effect of Cold Work and
`Heat Treatment on the R—phase to Austenite and Martensite to Austenite Transformation of a
`Nea—Equiatomic Ni—Ti Shape Memory Alloy”, Proceedings, International Conference on
`Martensitic Transformations, Monterey, California, 1993, pp. 917-922.
`[7] Miyazaki, 8., Engineering Aspects of Shape Memory Alloys, Butterworth-Heinemann
`(London, 1990), pp. 393—413.
`[8] ASTM F2005-05, “Standard Terminology for Nickel-Titanium Shape Memory Alloys”,
`ASTM,2005.
`[9] ASTM F2004~03, “Standard Test Method for Transformation Temperature of Nickel—
`Titanium Alloys by Thermal Analysis”, ASTM, 2003.
`[10] ASTM F2082-O3, “Standard Test Method for Determination of Transformation Temperature
`of Nickel-Titanium Shape Memory Alloys by Bend and Free Recovery”, ASTM, 2003.
`[11]Otsuka, K. and Wayman, C.M., “Mechanism of Shape Memory Effect and Superelasticity”,
`Shape Memory Materials, (Cambridge University Press 1998), pp.27—48.
`[12] Olson, GB. and Cohen, M., Scripta Met, v01. 9, 1975, pp.1247.
`
`l ] I
`
`154
`
`Edwards Exhibit 1022, p. 12
`
`