JOURNAL OF ENDODONTICS
`Copyright © 2002 by The American Association of Endodontists
`
`Printed in U.S.A.
`VoL. 28, No. 10, OcTOBER 2002
`
`Fatigue and Mechanical Properties of Nickel(cid:173)
`Titanium Endodontic Instruments
`
`Gregoire Kuhn1 AMU, and Laurence Jordan, MCU-PM
`
`Shape memory alloys are increasingly used in super(cid:173)
`elastic conditions under complex cyclic deformation
`situations. In these applications, it is very difficult to
`predict the service life based on the theoretical law.
`In the present work, fatigue properties of NiTi engine(cid:173)
`driven rotary files have been characterized by using
`differential scanning calorimetry (DSC) and mechan(cid:173)
`ical testing (bending). The DSC technique was used
`to measure precise transformation. The degree of
`deformation by bending was studied with combined
`DSC and mechanical property measurements. In
`these cold-worked files, the high dislocation density
`influences the reorientation processes and the crack
`growth. Some thermal treatments are involved in
`promoting some changes in the mechanical proper(cid:173)
`ties and transformation characteristics. Annealing
`around 400°C shows good results; the recovery al(cid:173)
`lows a compromise between an adequate density for
`the R-Phase germination and a low density to limit
`the brittleness of these instruments. In clinical usage,
`it is important to consider different canal shapes. It
`could be proposed that only few cycles of use is safe
`for very curved canals but to follow the manufactur(cid:173)
`er's advise for straight canals.
`
`In endodontic treatments, the risk with traditional files (stainless
`steel) is plastic deformation and fracture. Consequently, nickel(cid:173)
`titanium (NiTi) instruments with pseudo-elasticity properties
`(shape memory effect and superelasticity) have been introduced to
`avoid or to limit the failure risk. However, cyclic deformation
`during endodontic treatment changes the mechanical behavior of
`NiTi alloys and finally leads to fatigue failure.
`The superelasticity (SE) nature ofNiTi has been attributed to a
`reversible austenite to martensite transformation. It is believed
`austenite is transformed to martensite during loading and reverts
`back to austenite when unloaded. The transformation is reversible
`during clinical use, because SE alloys have a transition temperature
`range (TTR) lower than mouth temperature. The TTR of NiTi is
`effected by the chemical composition, method of fabrication, and
`heat treatment of the alloy (I). Sometimes, the direct transforma-
`
`tion from austenitic to martensitic NiTi includes an intermediate
`structure, called R-phase. It is important to have knowledge of the
`relationships between Austenite, R-phase, and Martensite transfor(cid:173)
`mation sequences on cooling and heating. On cooling, we can
`observe: Austenite ~ R-phase ~Martensite (direct transforma(cid:173)
`tion) and on heating: Martensite~ R-phase ~Austenite (reverse
`transformation). However, due to the large differences between the
`hysteresis of the martensite and the R-phase transformations, in
`some cases, the transformation sequence on heating is Martensite
`~Austenite directly. The transformation A~ R-phase shows the
`same properties (superelasticity and shape memory effect) because
`of the quasi-martensitic nature of this transformation. Young's
`Modulus of the R-phase is lower than that of Martensite, and thus,
`an instrument with the R-phase transformation would be more
`flexible.
`The mechanical properties and various phase transformation tem(cid:173)
`peratures ofNiTi shape memory alloy are known to be very dependent
`on thermo-mechanical processing. To use these alloys in various
`applications, proper control and understanding of the effects of
`thermo-mechanical processing parameters are very essential. We used
`some thermal treatments that are involved in promoting some changes
`in the mechanical properties and transformation characteristics. These
`properties can be modified by high dislocation density and/or fme
`dispersion of particles. A variety of irreversible phenomena associated
`with dislocations, precipitates, and residual stresses from previous
`cold work and/or thermal histories complicates the exploitation of
`superelastic (SE) alloys.
`The aim of this work is to show fatigue characteristics of
`superelastic NiTi, and subsequently, the effect of the process
`history on fracture life. This work is the continuation of our results
`based on microstructural investigations (scanning electron micros(cid:173)
`copy, X-rays diffraction, and microhardness) of nickel-titanium
`instruments (2). We investigated mechanical properties with em(cid:173)
`phasis on flexibility of endodontic instruments by the use of
`bending tests. This study will then be discussed using differential
`scanning calorimetry (DSC) results. DSC allows the identification
`of crystallographic phases at various temperatures (3).
`
`MATERIALS AND METHODS
`
`Materials
`
`The engine-driven, rotary instruments that were studied are
`produced by Maillefer (ProFile, Ballaigues, Switzerland) and by
`
`716
`
`GOLD STANDARD EXHffiiT 2021
`US ENDODONTICS v. GOLD STANDARD
`CASE IPR2015-01476
`
`

`
`Vol. 28, No. 10, October 2002
`
`Mechanical Properties of NiTi Instruments
`
`717
`
`TABLE 1. NiTi specimens
`
`Conicity-Diameter
`
`0.02/30
`0.04/20
`0.04/30
`0.06/20
`0.06/30
`
`Hero
`}
`
`}
`
`}
`
`ProFile
`
`} ✲ Œ
`}
`} 夽
`}
`
`} ⫽ new instrument, active part; ✲ ⫽ new instrument, inactive part; 夽 ⫽ used
`instrument, active part; Œ ⫽ instrument after thermal treatments, active part.
`
`in many geometrical
`Micro-Mega (Hero, Besanc¸on, France)
`shapes (Table 1). The studied files have a 25-mm length, a taper
`ranging between 0.04 and 0.06 mm per mm length, and sizes 20 to
`40, representing the diameter of the tip base of the file, given in
`hundredth of millimeter.
`Specimens were cut to separate working or active part of the file
`from the inactive part using a low-speed diamond saw. Several
`samples were chosen: new instruments and instruments that have
`been used in clinical conditions (12 or 18 root canals and approx-
`imately 5 or 6 sterilizations).
`
`Thermal Treatments
`
`Different thermal treatments were investigated. The heat treat-
`ments consisted of anneals at 350°C, 400°C, 450°C, 510°C, 600°C,
`and 700°C in salt baths for 10 min and at 600°C and 700°C for 15
`min with the same process and subsequent water quench in all
`cases.
`Machining process promotes a high density of defects; the alloy
`is work-hardened. We used some thermal treatments, which are
`involved in promoting changes in the mechanical properties and
`transformation characteristics. The influence of different heat treat-
`ments on NiTi alloys was investigated by DSC measurements and
`bending tests.
`
`Methodologies
`
`DIFFERENTIAL SCANNING CALORIMETRY
`
`DSC testing is one of the many test methods used to measure
`transformation temperatures of NiTi alloys. DSC testing is a ther-
`mal method that measures the change in heat flow, which is
`associated with the martensitic and austenitic phase transforma-
`tions through a controlled cooling/heating cycle. In the DSC pro-
`cedure, the differential heat flow required to heat or cool the
`experimental and reference samples at the same scanning rate is
`recorded as a function of temperature to yield the spectrum or
`thermogram. The start and finish temperatures of each phase trans-
`formation were determined from tangent lines where the DSC
`curve deviates from the adjacent baselines.
`The transformation temperatures were determined by DSC
`(Mettler 30/TA 4000). The file specimens were carefully cut with
`an oil-cooled, diamond-embedded saw. Considerable care was
`taken in cutting the samples so that minimal heat and stress would
`be generated. Straight segments of 5-mm length were cut from
`each file. Specimens were placed in aluminum pans with a nitrogen
`gas flow environment to prevent condensation of water vapor on
`the NiTi specimens. Another empty aluminum pan served as an
`inert reference. The weight of the samples for DSC measurements
`was 18 mg; heating and cooling rates were 5°C/min. The thermo-
`
`FIG 1. DSC thermograms obtained for different conicity.
`
`grams were carried out in temperature intervals of ⫹60°C and
`⫺120°C. During measurements, the samples were quickly heated
`to 60°C and then cooled to approximately ⫺120°C at a constant
`cooling rate (5°C/min). When the low temperature was reached,
`the specimen was heated again to 60°C at the same rate.
`
`Bending Tests
`
`To perform bending tests, we used a bending testing machine.
`All the instruments were loaded with the same deformation, and
`the forces corresponding were calculated by the cell (100 N). The
`loading and the unloading were performed in the same conditions.
`We tried with our machine to reproduce the bending of files that
`occurs in clinical situations. New instruments, instruments used in
`the clinic, and instruments that have been heat-treated were in-
`cluded in this mechanical test. We obtained information about the
`elastic behavior (flexibility) of files and about heat treatments and
`clinical use. The results are discussed only in a qualitative analysis
`and not a quantitative analysis because of the shape of the instru-
`ments (range and machining design), which prevents any calcula-
`tion.
`
`RESULTS
`
`DSC
`
`The transformation process can easily be recorded by measuring
`the transformation latent heat released/absorbed to/from the sur-
`roundings. Figure 1 shows the DSC curves for different conicities
`of Hero instruments. DSC curves show one-step distinct transfor-
`mation during cooling. The transformation product of the first
`exothermic peak at higher temperature corresponding to the trans-
`formation from austenite (A) to R-phase (R), whereas that of the
`second exothermic peak at lower temperature is difficult to identify
`and is dissimulated in the baseline of the thermogram. Two-step
`endothermic transformation occurs during heating (M 3 R and R
`3 A).
`The peaks are better and better defined with increasing conicity.
`The ProFile samples show the same characteristics, but all peaks
`are less defined. In the oral environment, the specimen are com-
`pletely austenitic or a mixture of austenite and R-phase.
`
`

`
`718
`
`Kuhn and Jordan
`
`Journal of Endodontics
`
`TABLE 2. Transformation temperatures
`
`Direct transformation TTR
`(°C)
`
`Reverse transformation
`(°C)
`
`Hero
`ProFile
`
`20
`35
`
`23
`39
`
`FIG 2. DSC curves obtained from different part of the instrument.
`
`FIG 3. DSC thermograms showing new ProFiles and ProFiles that
`have been used in clinical conditions.
`
`Table 2 summarizes the main transformation temperatures for
`each file.
`DSC cooling and heating curves in Fig. 2 show peaks associated
`with the A 7 R transformation but poorly resolved peaks; we can
`observe two different TTR for the active part (35°C) and inactive
`part (41°C).
`Figure 3 shows the DSC curves for new ProFiles and ProFiles
`that have been used in clinical conditions (12 root canals and
`approximately 10 sterilizations). In the case of the new files, one
`peak is obtained on cooling and on heating (A 7 R). However,
`when the file is used, the same peaks are observed but the TTR
`shifts to a lower temperature: approximately 15°C lower. The R 7
`M transformation does not appear.
`After heat treatments, the samples show two different types of
`transformation courses during DSC measurements. The peaks in-
`tensity and position change with varying annealing conditions.
`
`FIG 4. (A) DSC curves after heat treatments below 600°C. (B) DSC
`curves after heat treatments above 600°C.
`
`After heat treatments below 510°C (Fig. 4A), two peaks are found
`during cooling process (A 3 R, R 3 M) and two peaks during
`heating (M 3 R, R 3 A). Above 510°C, thermal treatments yield
`a transformation behavior with one peak during cooling and during
`heating (Fig. 4B). When the annealing temperature is above the
`recrystallization temperature (600°C), only one exothermic peak
`can be found during cooling (A 3 M). We have the direct and
`reverse A 7 M transformations; the R-phase transformation does
`not exist. This thermal treatment shifts the martensite transforma-
`tion to lower temperatures. For these different thermal treatments,
`DSC thermograms show TTR evolution (Table 3).
`
`Bending Tests
`
`At first, and until 3 mm of strain, only the tip of the instrument
`is bent. Then, between 3 and 6 mm, the curvature is in the middle
`of the file. Finally, above 6 mm, the part that has the maximum
`cross-sectional area near the handle becomes deformed in turn.
`As can be seen from the curves, the samples deformed at room
`temperature recover their original state, indicating that the trans-
`formation temperature is close to room temperature. Specimens
`seem to exhibit significant plateau-stress drop or stress relaxation
`during superelastic cycling.
`
`

`
`Vol. 28, No. 10, October 2002
`
`Mechanical Properties of NiTi Instruments
`
`719
`
`Before Heat Treatment
`
`TTR
`
`35°C
`
`400°C
`
`40°C
`
`510°C
`
`22°C
`
`600°C
`⫺42°C
`
`700°C
`⫺46°C
`
`TABLE 3. TTR evolution after heat treatments
`
`stress reversal probably leads to defect formation. Dislocations
`present in the matrix influence the mechanical properties; internal
`stresses are a negative factor to the mobility of martensite inter-
`faces (6). Moreover, DSC thermograms (Fig. 3) and bending
`curves (Fig. 5) of used instruments (abrupt curvatures) show an
`increased density of dislocations. Indeed, the shift of the TTR to
`lower temperature impedes the phase transformation; instruments
`become stiff. It seems that the abruptness of the curvature of the
`root canals (stress field) is the essential parameter and that the
`numbers of root canals treated is less important for the increase of
`brittleness. Manufacturers advise not to use each file on more than
`10 to 12 root canals. The difference between files that have been
`stressed in curved canals compared with straight canals can be seen
`in terms of different defect densities created by martensite reori-
`entation. In the curved canals, the stress to induce martensite will
`be high. The build-up of stress concentrations could be prevented
`from forming cracks, because martensite will be retained to ac-
`commodate this stress. But if the density of defects is high (in case
`of work hardening), the reorientation of variants in the stress field
`or its reversion and reformation is not possible and cracks appear
`and grow.
`In Figs. 3 and 5, specimen behavior is probably a consequence
`of dislocations and lattice defects generated by high stress levels.
`When lattice defects are created, they will restrict the easy move-
`ment of martensite, i.e. they decrease the mobility of martensite
`interfaces.
`It was found that martensite transformation propagated in steel
`retarded crack growth, and it was proposed that this could result
`from internal compressive stresses induced by a positive volume
`change near the crack tip. However, in NiTi the volume change is
`small and negative and thus causes a negligible effect of stress-
`induced martensite near a growing crack (7).
`According to Yang (8), the fracture life apparently obeys the
`following relationship if high strain (curved canal) and low strain
`are considered separately according to different deformation mech-
`anisms:
`
`⌬⑀ 䡠 N␤ ⫽ C
`
`where C and ␤ are constants, ⌬␧ is the applied strain amplitude,
`and N is the number of cycles to fractures. For an optimization of
`fatigue resistance in the specific range of reversible deformation, it
`is necessary to pay attention to the shape of canals. The fatigue
`behavior of superelastic files is highly dependent on strain level
`and heat treatment conditions.
`Heat treatments are known to influence mechanical properties
`and various phase transformation temperatures of NiTi shape
`memory alloys. Two annealing temperature ranges can be distin-
`guished. In the first case, annealing temperatures approximately
`600°C (recovery) show two-step transformations (A 7 R 7 M);
`in the second case, when annealing temperatures are above 600°C
`(recrystallization), we can observe a direct martensitic transforma-
`tion. The annihilation of dislocations by recovery and recrystalli-
`zation, or the beginning of dissolution of precipitates, are of special
`importance both for the structural properties and the functional
`properties (e.g. transformation temperatures). It can be supposed
`that voids that were quenched into the microstructure will behave
`
`FIG 5. Bending curves showing new ProFiles and ProFiles that have
`been used in clinical conditions.
`
`We cannot notice any difference between a new instrument and
`a file that has been used in straight canals (Fig. 5). But when the
`abruptness of canal curvature increases, the stiffness of the used
`file increases after each use.
`Figure 6 (A and B) demonstrate that the annealing conditions
`strongly affect the stress-strain behavior. For heat treatments below
`recrystallization temperature (Fig. 6A), the specimens generally
`show an increased flexibility. On the other hand, results show that
`after annealing at a temperature above recrystallization, the stiff-
`ness of the instruments increases (Fig. 6B).
`
`DISCUSSION
`
`Results previously presented with XRD and microhardness (2)
`and our results with DSC show that new specimens are signifi-
`cantly work-hardened. The microhardness measure is twice that of
`a fully recrystallized sample (4). Moreover, machining marks and
`cracks on the surface observed by scanning electron microscopy
`contribute largely to fatigue failure by a crack propagation process.
`The crack nucleation stage is improved by the high density of
`surface defects. Manufacture of NiTi alloys by machining into
`endodontic instruments promotes work hardening and creates sur-
`face defects. These first observations (microstructure and surface
`defects) explain unexpected fractures reported in the literature (5).
`A variety of irreversible phenomena associated with disloca-
`tions, precipitates, ordering effects, and residual stresses from
`previous cold work and/or thermal history complicates the use of
`shape memory alloys. The substantial variance in the thermograms
`of active and inactive parts results from the manufacturing pro-
`cesses. The active part is more affected by machining compared
`with the inactive part (Fig. 3). When the deformation temperature
`(oral temperature) lies near Ms on initial loading, a martensite and
`subsequently R-phase are stress induced and the latter is stable
`throughout the rest of the test. Those martensite (or R-phase)
`variants will be selectively produced, which gives maximum strain
`in the direction of the applied stress and their reorientation on
`
`

`
`720
`
`Kuhn and Jordan
`
`Journal of Endodontics
`
`strongly depend on test temperature (9); such strong temperature
`dependence is unique and cannot be found in the standard data of
`ordinary materials.
`Various degradation modes occur, including a shift in transfor-
`mation temperature, a reduction in the available strain, and of
`course, fracture. With regard to the fatigue behavior in the super-
`elastic temperature range, fracture is a critical failure mode for
`many applications.
`In these cold-worked files, the high dislocation density influ-
`ences the reorientation processes and the crack growth. The in-
`struments become brittle. As to superelasticity, with cycling reori-
`entation of the martensite under stress leads to gradual defect
`accumulation, and it might be expected that these dislocations are
`generated at the interface between different martensite colonies. In
`clinical conditions, the curvature of canals distorts the endodontic
`instruments; cyclic fatigue is caused by repeated tensile-compres-
`sive stress. The maximum of this stress is in the surface of the
`curve. Crack nucleation and propagation stages appear mostly on
`the half-part of the instrument, which is in tension (outside of the
`curve).
`Some suggestions could be proposed to improve the lifetime of
`endodontic files; these include applying thermal treatments at
`approximately 400°C (recovery) before machining to decrease the
`work-hardening of the alloy, choosing machining conditions
`adapted to this NiTi shape memory alloy, and electropolishing by
`the manufacturer to reduce the machining damage on the file
`surface. For an optimization of the fatigue resistance in the specific
`range of reversible deformation, it is necessary to pay attention to
`shape of the canal. Only a few cycles of use for very curved canals
`may be best, and following the manufacturer’s advice may be best
`for straight canals. Therefore, it is very important to understand the
`fatigue characteristics of NiTi alloys to use these functions in
`various types of applications.
`
`The authors are grateful to Dr. France Dalle for reviewing the manuscript.
`
`Drs. Kuhn and Jordan are affiliated with the Faculty of Dentistry, University
`Denis Diderot, Paris, France. Address requests for reprints to Gre´ goire Kuhn,
`Laboratoire de Me´ tallurgie Structurale, ENSCP, 11 Rue Pierre et Marie Curie,
`75231 Paris Cedex 05, France.
`
`References
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`Met 1979;27:137– 44.
`8. Yang J. Fatigue characterization of superelastic nitinol. SMST. Pro-
`ceedings of the Second International Conference on Shape Memory and
`Superelastic Technologies. 1997, Pacific Grove, CA.
`9. Kim YN, Miyazaki S. Fatigue properties of Ti-50.9 at% Ni shape memory
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`
`FIG 6. (A) Bending curves for various annealing conditions. For heat
`treatments below 600°C, the specimens show an increased flexibil-
`ity. (B) Bending curves for various annealing conditions. For heat
`treatments above 600°C, the stiffness of the instruments increases.
`
`differently depending on whether there are many dislocations
`during annealing. Due to the thermal activation, the voids will
`move into the fields of compressive stress (inside the dislocations)
`and reduce the number of nucleation sites in a microstructure with
`higher dislocation density. With beginning recrystallization and
`dissolution of the particles after further annealing, no more R-
`phase transformation can be detected. With the beginning recrys-
`tallization (600°C and 700°C) (Figs. 4B and 6B), R-phase trans-
`formation cannot be detected. It can be presumed that
`the
`decreasing dislocation density and any precipitation stress fields
`are not able to initiate the R-phase transformation. The dissolution
`of Ni-rich precipitates increases Ni-content in the matrix and shifts
`the TTR to a lower temperature. Thus, the stiffness is much more
`important. For clinical applications, these heat treatments are not
`required.
`The shape memory alloy files are expected to show unique
`fatigue characteristics different from the usual materials because of
`their deformation behavior associated with the martensitic trans-
`formation. In fact, fatigue crack propagation rate and fatigue life