reversible magnetic field induced strain in ni2mnga-polymer-composites

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Reversible Magnetic Field Induced Strainin Ni2MnGa-Polymer-Composites**

By Sandra Kauffmann-Weiss,* Nils Scheerbaum, Jian Liu, Hansjorg Klauss, Ludwig Schultz,Edith Mader, Rudiger Habler, Gert Heinrich and Oliver Gutfleisch

Composite materials consisting of magnetic shape memory alloy particles and a polymer matrix combinethe advantages of both material classes: the high achievable magnetic field induced strain (MFIS) of 6% ofNi-Mn-Ga with a ductile matrix. Engineering the particle-matrix interface as well as matching stiffnessof polymer matrix is of importance for achieving high reversible MFIS to use this material as actuator ordamper. We investigated those properties for Ni50.9Mn27.1Ga22.0 and Ni50.3Mn24.6Ga25.1 polymercomposites. Particles were produced by gently crushing melt-extracted and subsequently annealedfibres. At room temperature, the Ni50.9Mn27.1Ga22.0 particles exhibit a 5M martensitic structure, whilethe Ni50.3Mn24.6Ga25.1 particles are austenitic. These particles were embedded into the polymer, either astiff epoxy resin or a soft polyurethane. In response to an external appliedmagnetic field, the particles tendto relocate within the polyurethane due to its very low Young’s modulus and magnetostatic interactionbetween particles. Slightly stiffer polymer matrices are advantageous for achieving controllable MFIS. InNi50.9Mn27.1Ga22.0 epoxy composites, a MFIS of 0.1% was observed and was resettable by rotating themagnetic field by 908. Furthermore, single fibre pull-out tests indicated significant improvements of theinterfacial properties when using silane coupling agent treated fibres.

Magnetic shape memory (MSM) alloys are a new class of

materials for sensor and actuator applications. Ni-Mn-Ga

single crystals show large magnetic-field induced strain

(MFIS) in moderate magnetic fields below 1 T caused by

reorientation of martensitic variants by twin boundary

motion. The maximum possible strain e¼ 1-(c/a) is given

by the lattice parameters a and c of the martensite unit cell

resulting in e¼ 6% for 5M martensite and e¼ 11% for 7M

martensite.[1–3] However, the main disadvantage of single

[*] S. Kauffmann-Weiss, Dr. N. Scheerbaum, Dr. J. Liu, H. Klauss,Prof. L. Schultz, Dr. O. GutfleischIFW Dresden, Institute for Metallic Materials,PO Box 270116, 01171 Dresden, (Germany)E-mail: s.weiss@ifw-dresden.de

Prof. E. Mader, Dr. R. Haßler, Prof. G. HeinrichLeibniz-Institut fur Polymerforschung Dresden e.V.,Hohe Straße 6, 01069 Dresden, (Germany)

S. Kauffmann-Weiss, Prof. L. Schultz, Prof. E. Mader,Prof. G. HeinrichTU Dresden, Institute for Materials Science,D-01069 Dresden, (Germany)

[**] This work is supported by DFG SPP 1239. Experimentalassistance by the Research Technology Department of theIFW Dresden is gratefully acknowledged.

20 wileyonlinelibrary.com � 2012 WILEY-VCH Verlag GmbH & Co.

crystals is their high brittleness. One possible solution is

the synthesis of textured polycrystalline Ni-Mn-Ga for

which a MFIS of 1% has been reported.[4] Due to constraints

by grain boundaries, the twinning stress stwin for moving

twin boundaries is still much higher for polycrystals

(stwin> 15MPa) compared to that in single crystals

(stwin< 2MPa). A strong texture, large grains and a mechan-

ical training can decrease stwin.[4,5] The main disadvantage of

single- and polycrystals is their complex preparation.

An alternative to single- and polycrystals are Ni-Mn-Ga

polymer composites.[6–8] Single crystalline particles

embedded in a stiffness-matched polymer matrix allow

overcoming the disadvantages. A thin polymer film between

the single crystalline particles allows the particles to strain

und reduces stwin. Another advantage of composites is the

decrease of eddy currents due to a non-conducting polymer

matrix, which enables high frequency applications, and the

simple preparation of textured bulk materials.

In previous works, we showed that melt-extracted and

subsequently annealed fibres exhibit a bamboo-like grain

structure and an MFIS of 1%.[6] However, only small amount

of grains in the fibre are active. Different crystallographic

orientations of grains as well as still existing grain boundaries

hinder activation of the whole fibre. A separation of fibres into

nearly single crystalline particles by breaking them along

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grain boundaries should increaseMFIS. Composites with a stiff

polymermatrix can dissipate a large amount of input energy by

stress induced twin boundary motion within the particles

and therefore can be used as dampers.[7,8] The application as

actuators requiresmagnetically induced twin boundarymotion

and thus a polymer matrix with a lower stiffness is necessary.

Alignment of the Ni-Mn-Ga particles with their easy magne-

tisation axis parallel to each other can be achieved by applying

a magnetic field during curing of polymer matrix.[8] In case of

the cured composite, this results in an easy magnetisation

behaviour for the designated direction and harder magnetisa-

tion behaviour for the perpendicular directions.

For the transfer of MFIS from the particles to the entire

composite, a good adhesion of particle and polymer matrix is

important. As the strain within the composite is transferred

partly by shear strain of the polymer, a poor adhesion between

particles and polymer facilitates particle rotation. Further, the

polymer would dissipate the strain of the particles, too.

In this work, the interfacial properties and the influence of

the polymer matrix stiffness in magnetic Ni-Mn-Ga polymer

matrix composites are investigated. It will be shown, that

those parameters are crucial for obtaining reversible MFIS.

1. Experimental

1.1 Preparation and Analysis Methods

To compare the magnetic and elastic properties of

composites, particles with different crystallographic struc-

tures at room temperatures were used. The structure of the

Ni-Mn-Ga system has a strong dependency on the composi-

tion, i.e. variations in the range of 0.1 at% may result in

different structures.[11] To investigated structure and marten-

sitic transformation temperatures DSC, suszeptometer,

SQUID and XRD measurements were performed.

Ni50.9Mn27.1Ga22.0 and Ni50.3Mn24.6Ga25.1 (at%, determined

by inductively coupled plasma optical emission spectrometry

ICP-OES) fibres were prepared by crucible melt extraction.

Annealing of fibres was done under argon atmosphere at

1100 8C for 2 h, followed by slow cooling to room temperature

(furnace cooling till 700 8C). Particles were produced by

gently crushing annealed fibres.[8] Thematrixmaterials were a

two-component epoxy resin (epple 07170, E. Epple & Co

Fig. 1. a) Quasistatic fibre pull-out arrangement with load cell (1), fibre (2) and polymer matrix (3). Opticalstrain measurement equipment: relative orientation of sensor (4), sample (5), aluminium pad (6) and magneticfield for b) parallel and c) perpendicular strain measurements.

GmbH, hardener:resin ratio¼ 15:85) and a

two-component polyurethane resin (UR 5048,

Electrolube, hardener:resin ratio¼ 7:93). The

composites were prepared by mixing the

particles with the liquid polymer (filling

degree: 50 vol.% particles) in a mould of

10� 10� 40 mm3, followed by vacuum treat-

ment to remove remaining air and subse-

quently curing on air in a magnetic field of 0.5

T. The cured composite sampleswere cutwith

a standard diamond saw to 10� 10� 10mm3.

X-ray diffraction (XRD) was performed at

room temperature with Philips X’Pert using

cobalt Ka radiation. The martensite-austenite

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transformation temperatures and the Curie temperature were

determined using a differential scanning calorimeter (DSC,

Perkin Elmer Pyris 1, heating and cooling rate of 10 K �min�1),

an a.c. susceptometer (6 kHz, heating and cooling rate

� 1K �min�1) and a SQUID magnetometer (Quantum Design

MPMS-5S, constant field of 10 mT for magnetization vs.

temperature loops). Scanning electronmicroscopy (LEOGemini

1530 FEG-SEM) was used to analyse the microstructure of

particles and the pulled-out fibres regarding the failure

mechanisms. Grain/twin boundaries and texture information

were obtained by electron backscatter diffraction (EBSD,

HKL Channel 5). The Young’s modulus of polymers was

determined in uniaxial compression (cubic, sample size about

10� 10� 10mm3) and in tensile tests (dumbbell shaped speci-

mens: ‘‘S2’’ shouldered test bar, measuring length 20mm,

Instron Instruments 8562). With dynamic-mechanical analysis

(DMA) the storage and loss modulus were tested (Analysator

Q800, TA Instruments, single cantilever, heating rate 3 K �min�1,

frequency 1Hz, amplitude 20mm, free clamping length

17.5mm). With this the Young’s modulus can be calculated.

1.2 Optical Strain Measurement Equipment

For the MFIS measurements the composite sample was

connected with glue to the sample holder on one side and to a

reflecting aluminium pad on the opposite side. The alumi-

nium pad reflected the light from the confocal displacement

sensor (micro epsilon, controller IFC2401, sensor IFS2402/

90-4), which is used fro strain measurements. The sample

holder could be rotated with respect to the applied magnetic

field direction. Figure 1b and c show the parallel and

perpendicular alignment between strain measurement and

applied magnetic field direction. The magnetic field was

applied by standard electromagnet. The resolution of the

equipment was < 1mm and a maximal magnetic field of �1.8

T could be applied at room temperature.

1.3 Quasistatic Pull-Out Test

In order to investigate the particle-matrix-interface proper-

ties, quasi-static pull-out tests were performed. About 15

single fibre composites were prepared without and with

epoxysilane coupling agent (GLYMO, Dynasylan, Evonik

Industries) treated fibres. These composites were made by

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Fig. 2. a) DSC (top) and suszeptometer (bottom) analyses of the at room temperature martensitic (solid) and austenitic (dashed) Ni-Mn-Ga particles. b) SQUID magnetometer (top)analysis and XRD pattern at room temperature (bottom). For the austenitic particles the three austenitic peaks and for the martensitic particles the typical 5M modulated martensitestructure peaks were present.[12]

using separate sample preparation equipment, as described

elsewhere.[9] Single fibres were embedded into the polymer

matrix with a pre-selected embedding length (Figure 1a).

Embedded lengths le were chosen in the range of 150mm to

475mm depending on the fibre diameter df (le � 5df).[10]

The fibre diameter was measured by optical microscopy. The

modified single fibre composites were prepared in the same

way with the exception that the fibres were treated with

Glymo before embedding. All single fibre composites were

cured at room temperature for 3 weeks in exsiccator. The

Ni-Mn-Ga fibres were pulled out of the polymer matrix at

equal ambient conditions with a oading rate of 0.01mm � s�1.

During pull out, the force-displacement curve was measured.

With maximum force Fmax, embedded length and fibre

diameter, the apparent shear strength tapp can be calculated

with the following equation

tapp ¼Fmax

pdf le(1)

At Fmax a complete interface debonding of fibre and matrix

occurs (adhesion and friction).[10] This simple reduction

allows only a comparison with different fibre sizes and

models.

2. Results and Discussion

Fig. 3. Theoretical stress-strain curve of Ni-Mn-Ga and the simplification for calculation.

2.1 Ni-Mn-Ga Particles

Ni-Mn-Ga particles were obtained by

gently grinding melt-extracted fibres. These

fibres were about 40 to 100mm in diameter

and several millimetres in length. After

annealing, the fibres exhibited a bamboo-like

grain structure with grains as large as the

fibre diameter. By gently grinding, the fibres

broke preferably and easily along grain

boundaries.[8] About 50% of particles were

single- or oligo-crystalline. In order to

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investigate austenite-polymer-composites, ferromagnetic

Ni50.3Mn24.6Ga25.1 particles (dashed curves in DSC, suscept-

ometer, and SQUID measurements, Figure 2) with martensi-

tic-austenitic transformation temperature below room tem-

perature (� 200K) were used. The XRD pattern in Figure 2

occupies the austenitic structure. For the MFIS measurements

on martensite-polymer-composites, Ni50.9Mn27.1Ga22.0 parti-

cles (solid curves in DSC, susceptometer and SQUID

measurements, Figure 2) were chosen. For these particles

the martensitic-austenite transformation took place between

314 and 338K and the 5M-martensite was present at room

temperature. Due to the martensitic structure in the XRD

pattern in Figure 2 the peak splitting was clearly seen.

2.2 Polymer Selection

The elastic polymer film between particles should be soft

enough to allow the particles to strain, because the maximal

magnetic stress tmax,M applicable to induce twin boundary

motion (stwin � 2MPa) is in the range of a few MPa (Figure 3)

and given by [13]

tmax;M ¼ K

sfor H � HS (2)

The magnetic anisotropy constant for a 5M martensite is

K¼ 1.6� 105 J �m�3 and the twinning shear s is defined by the

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Fig. 4. Calculated strain of Ni-Mn-Ga polymer composites as a function of volumefraction of Ni-Mn-Ga and polymer elastic modulus EPolymer. The expected strain is of3.1 and 0.6, for the polyurethane and epoxy matrix respectively.

distance of twin planes (dt) and the burgers vector (b) with

s¼ dt/b¼ 0.062.[14] For 5M martensite tmax,M is about

� 2.6MPa. For a very stiff polymer the internal-stress between

particle and matrix is in the order or larger than tmax,M. So the

polymer matrix blocks particles and hinder the possible MFIS

of those. Therefore, the stiffness of the polymer matrix is of

vital importance for MFIS in composites. One approach to

calculate the expected composite strain in dependence of

matrix stiffness is the stress-free strain method.[15] The

stress-free strain e� is given by following equation:

"� ¼ "T þ Sþ S�� �

SP � SMð Þ "TM � "TP

� �(3)

Using the compliance tensors of matrix SM and particles SP,

the effective compliance tensor S�, the stress free strains of

matrix eM and particles eP, the average stress free strain "T and

the average compliance tensor S.

With the assumption that the particles are homogenously

distributed in the polymer matrix and the elastic properties of

particles, matrix and composite are isotropic, an alternative

expression for the stress-free strain e� can be obtained by

equation 4.[15]

"� ¼ "T þ VPVMKPM1

KP� 1

KM

� �"TM � "TP� �

(4)

KPM ¼ VM

KPþ VP

KMþ 3

4GM

� ��1

(5)

Thereby e� is defined by the volume fraction of the particles

(VP) and the matrix (VM), the elastic constants of the particles

(bulk modulus KP) and the matrix (bulk modulus KM

and shear modulus GM) and the MFIS of particles ("TM) and

matrix ("TP).

In Figure 3 a theoretical stress strain curve for a Ni-Mn-Ga

single crystal is shown. For a very small strain the deformation

behaviour can be described with the elastic Young’s modulus

(Eelastic> 2GPa). In the range of stwin the twin boundary is

moving and a strain (MFIS) of 6% is observed. However, for

calculations the stress-strain behaviour is simplified. Since no

MFIS occurs in the initial range the elastic Young’s modulus is

not regarded. To define the elastic constants for the equation 4

only the range for twin boundary motion is used (inset in

Figure 3). With applying amagnetic field or mechanical stress,

the 5M martensitic particles show MFIS up to "TP ¼ 6%,

while the polymer matrix does not ("TM ¼ 0%, Poisson ratio

nM¼ 0.33). With the assumption, that the twin boundaries

are relatively mobile and a constant stwin, the Young’s

modulus during twin boundary motion in Ni-Mn-Ga

particles is estimated to be Etwin¼EP¼ 1MPa (nP¼ 0.4).

Thus, one can calculate KM and GM by K¼E/(3(1� 2n))

and G¼E/(2(1þ n)).

With this simplification the composite strain (MFIS,

Figure 4) in dependence of the Young’s modulus of the

polymer matrix EM and the volume fraction of Ni-Mn-Ga

particles VP could be calculated. As expected the smaller the

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matrix modulus GM (therefore EM) and the larger the volume

fraction of particles VP the larger the composite strain e�. Thecalculated composite strain given here can be expected to

be an upper limit for the real magnetic field induced

composite strain.

For the experiment, two polymer matrices with different

stiffness were used: epoxy resin and polyurethane. The

Young’s modulus of those polymers had been investigated by

compression and tensile tests as well as DMA and different

strain rates of loading (Figure 5). In general, polyurethane had

a much smaller Young’s modulus (EM � 2MPa) than epoxy

resin (42 –858MPa). The large range of the Young’s modulus

of epoxy resin demonstrated the strong dependence on

loading type (DMA or compression) and strain rates, due to

relaxation of the polymer. After applying an external load, the

polymer fully relaxed into to the equilibrium during the

relaxation time. This happens through rearrangements of

main and side chain segments and rotation of end functions. A

further explanation is the long curing time of the epoxy resin.

After about 6 weeks, the epoxy resin was completely cured,

which is reflected in an increase of Young’s modulus till

858MPa. Longer curing time did not increase the Young’s

modulus further (results are not shown here).

Using equation 4, the determined Young’smodulus of both

polymers (results of compression tests after 3 weeks curing)

and a volume fraction of 50 vol% Ni-Mn-Ga particles yield to

an expected strain of 3.1 and 0.6, for the polyurethane and

epoxy matrix respectively.

2.3 MFIS in Composites

The Ni-Mn-Ga particles were mixed either with liquid

epoxy resin or polyurethane. The EBSD map of a martensi-

te-epoxy-composite in Figure 6a shows the distribution and

crystallographic orientation of Ni-Mn-Ga particles within the

epoxy matrix (white: not indexed by EBSD, represents the

epoxy matrix). The colour code represented the angular

deviation between crystallographic c-axis and the applied

magnetic field direction during curing (from 08¼ blue

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Fig. 5. Young’s modulus of polymers in dependence on a) strain rate (for tensile, compression tests and DMA) and b) curing time (Epoxy: strain rate for compression 8 � 10�4 s�1 andfor DMA 1 s�1; Polyurethane: strain rate for tensile 2 � 10�3 s�1).

Fig. 6. a) EBSD map (step size 0.5mm) of martensite-expoy-composite cured in a magnetic field of 0.5 T. Thecolour code represented the angular deviation of c-axis and applied magnetic field direction during curing (sheetnormal, deviation from 08¼ blue to 908¼ red). Black lines represented twin boundaries and white areas were notindexed (epoxy matrix). b) The [001]-pole figure correspond to the c-axis. A fibre texture along the z-directionwas seen in the pole figure.

(parallel) to 908¼ red (perpendicular)). Black lines in the inset

represented twin boundaries (defined by 868 h110i misor-

ientation between EBSD map pixels).[6] The particles, which

are homogeneously arranged within the epoxy matrix, were

single- or oligo-crystalline. Image analysis yielded a volume

fraction of Ni-Mn-Ga particles of about 50 vol. %. All grains

consisted of several twin variants. There was no correlation

between crystallographic orientation and particle shape.[6,16]

The pole figure (Figure 6b) showed a [001]-fibre texture

along the sheet normal (z-axis), being the axis along which

a magnetic field of 0.5 T was applied during curing

of composite. About 50% of embedded particles had a

misorientation of less than 358 between the crystallographic

c-axis and the sample’s z-axis. In the liquid polymer the

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c-axis can be oriented by rotation of particles

or by twin boundary motion during curing.

Texturing of Ni-Mn-Ga particles in compo-

sites by applying a magnetic field during

curing had been demonstrated in our pre-

vious work by magnetic measurements

and was here proven directly by the EBSD

analysis in.[8]

The calculation of the composite strain

showed that the softer the polymer matrix

the higher the expected magnetic field

induced composite strain (Figure 4). On the

other hand measurements on austenite-

polymer-composites showed that a certain

stiffness of the polymer matrix is necessary.

Austenitic Ni-Mn-Ga has a similar saturation

magnetisation Msat as the martensite

(difference in Msat about 10%), but shows

almost no MFIS (only conventional standard

magnetostriction < 0.01%).[13,17,18] In Figure 7

the deformation strain of austenitic

polymer composites is illustrate. The external

magnetic field was applied parallel and

perpendicular to the strain sensor. The

austenite-polyurethane-composite (curve 1) with the very

soft polyurethane matrix deformed in a magnetic field.

In contrast, the austenite-epoxy-composite (curve 2) with

the relatively stiff epoxy matrix did not deform in magnetic

fields. Due to the low Young’s modulus of the used

polyurethane, the magnetostatic interactions between

particles and the slightly inhomogeneous external magnetic

field gives rise to bending of the composite. Hence, soft

polymer matrices are advantageous in respect to maximum

MFIS and work output, but might not be suitable for

MSM-polymer-composites.[19]

Figure 8 shows the MFIS measured on martensi-

te-epoxy-composite parallel and perpendicular to the applied

magnetic field. TheMFISwas first measured perpendicular by

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Fig. 7. Strain measurements on austenite-epoxy-composite (curve 2) and austenite-polyurethane-composite (curve 1). The polyurethane-composite deformed in a magneticfield.

increasing the field up to 1 T and holding until the strain

remains constant. Subsequently the field was decreased down

to 0 T. In zero fields the samplewas rotated by 908 and the field

was increased again up to 1 T and held until the strain

remained constant. This procedure was repeated for several

cycles. For a martensite-epoxy-composite cured for 3 weeks,

the maximum strain was 0.04% for the first application of

magnetic field, but decreased with increasing number of

Fig. 8. a) MFIS measurement on a martensite-epoxy-composite after curing. (1) Perpendicb) MFIS measurement after magneto-thermal training. The first strain measurement perpenSubsequently the field was rotated by 908 for a following parallel strain measurement (not sperpendicular to the magnetic field yielded in 0.5% MFIS (2).

Table 1. Comparison of expected calculated (with the Young’s Modulus ECompression atte-epoxy-composite with 50 vol.% Ni-Mn-Ga before and after magneto-thermal training.

strain measurement Ecompression (8 � 10�4 s�1)

after curing 50 MPa

after magneto –thermal training 170 MPa

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magnetic field cycles (Figure 8a). Another phenomenon is that

the strain exhibited a pronounced time-dependence, due to

relaxation of the polymer matrix. The maximum strain was

achieved after a few minutes holding at the maximum

magnetic field.

Moreover, an enhanced strain of 0.1% can be achieved by a

subsequentmagneto-thermal training (Figure 8b). For this, the

composite was heated up to 80 8C and then cooled under a

magnetic field (parallel to the magnetic field axis during

curing) from the austenite state into the martensite state at

room temperature. The aim of this procedure was to induce a

single martensite variant state within particles. The measured

strain of 0.1% for a Ni2MnGa-epoxy-composite is in the range

of textured Terfenol-polymer-composites and piezoelectric

ceramics.[20,21] After decreasing the magnetic field down to

zero a residual strain of 0.5% is observed (Figure 8b, curve 1).

By applying a magnetic field in perpendicular direction this

strain is reversible resettable. However, the strain decreased

with a higher number of load cycles in the same way like in

untrained samples. And no residual strain occurs (Figure 8b,

curve 2).

There are several reasons for the measured strain being

smaller than the calculated strain (Table 1): As seen in the

EBSD map in Figure 6 the particles are not all single

crystalline. Out of it not all embedded particles are active.

The existing grain boundaries constrain twin boundary

motion. But even the single crystalline particles must not

necessarily be all active; they contain several twin boundaries,

ular to magnetic fiel; (2) Parallel to magnetic field; (3) Perpendicular to magnetic field.dicular to magnetic field yielded a MFIS of about 0.1% and a residual strain of 0.5% (1).hown here – strain was resettable). After rotating again the second strain measurement

8 � 10�4 s�1, assuming all particles are active) and measured strain of a martensi-

calculated ecomposite measured ecomposite

0.6% 0.04%

0.19% 0.1%

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Fig. 9. Comparison of average force-displacement curves of single fibre pull-out test foruntreated Ni-Mn-Ga fibre (dashed) and Glymo treated Ni-Mn-Ga fibre (solid) embeddedin epoxy resin.

Fig. 10. Results of single fibre pull-out test for untreated and Glymo treated Ni-Mn-Gafibres embedded in epoxy resin: maximum force as a function of embedded length.

which may hinder each others movements, especially in

untrained composites (difference of 0.56% between the

expected and measured MFIS). After training at 80 8C the

Young’s Modulus became higher up to 170MPa (Table 1).

However the measured MFIS was higher than in untrained

sample with the lower Young’s Modulus. During magne-

to-thermal training, cooling into the martensite with an

applied magnetic field, twin variants with cjjH are preferred.

This reduces the number of twin variants and twin boundaries

within grains and thus also the resulting constrains. And the

difference between measured and expected strain is smaller

(0.09%).

2.4 Particle Matrix Interface

Another crucial parameter for these composites is

the adhesion strength at the interface between particles

and polymer matrix. Increased adhesion strength avoids

rotation of particles and is important to transfer the MFIS

from the particle to the surrounding polymer matrix and with

that to the entire composite. In order to characterise the

interfacial adhesion strength, quasi-static fibre pull-out tests

were conducted. Figure 9 shows average force-displacement

curves for untreated and epoxysilane treated fibres in

comparison.

From each force-displacement curve, the maximal force

Fmax was determined. The determination of the typically used

debonding force Fd and therefore the debonding shear

strength td, the point where first debonding from matrix

start, was impossible, because the typical kink could not be

exactly located in the curves.[10] The fibres’ surfaces are not

smooth, therefore several small kinks are visible in the curves

up to the maximum force. Above Fd, fractional cracks and

friction in the debonding regions overlap with non-debonded

regions. Figure 10 shows the results for Fmax of the fibre

pull-out tests. For the treated fibres Fmax is higher than for

untreated fibres. At Fmax a complete interface debonding

between fibre and matrix occurs. With Fmax, the apparent

shear strength tapp of the untreated Ni-Mn-Ga fibres could be

determined to about 1 to 3MPa, depending on the embedded

length as well as on the fibre diameter (not shown here). tappwas in the range of tmax,M and therefore too low. The apparent

shear strength tapp is comprised of contributions from

Fig. 11. a) SEM image of an untreated, pulled out fibre. The clean fibre surface indicates a failure by adhesivefracture. b) SEM image of an epoxysilane treated, pulled out fibre. Epoxy rests on fibre surface indicate a failureby adhesive and cohesive fracture.

adhesion and friction and varies with the

embedded length as well as the fibre dia-

meter.

In composites, tapp should be larger than

tmax,M in order to avoid interface fracture

duringmagnetic field induced twin boundary

motion. The relatively low interfacial adhe-

sion strength confirmed another reason for

the discrepancy between calculated and

measured strains. The surface of fibres after

the pull out test is visualised by SEM

(Figure 11a). The clean fibre surface indicates

a complete interface failure by adhesive

fracture. The reason for the composites’

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decrease in magnetic field induced strain after each cycle is

likely to be a gradual fracture of interfaces between particles

and matrix.

A significant improvement of the interface stability could

be achieved by treating fibres with the epoxysilane coupling

agent Glymo before mixing with the polymer matrix.

The Glymo-treated fibres show apparent shear strengths tapp

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S. Kauffmann-Weiss et al./Reversible magnetic field induced strain

above in the range of 3 to 6MPa, which is significantly higher

than that for the untreated fibres. The remaining epoxy matrix

on the fracture surface of the epoxysilane treated fibre is

visible after the pull out test (Figure 11b). This indicates at

least a partial failure by cohesive fracture. A cohesive fracture

marks the improved interfacial adhesion strength.[22]

3. Conclusions

Martensitic Ni50.9Mn27.1Ga22.0 and austenitic Ni51.4Mn26.9Ga21.7 fibres were prepared by crucible melt extraction follow

by annealing at 1100 8C for 2 h. Composites were prepared by

embedding small single- and oligo-crystalline particles,

obtained by gently crushed annealed Ni-Mn-Ga fibres, in

two different polymer matrices. The epoxy resin (E¼ 42MPa

to E¼ 858MPa) is much stiffer than the polyurethane

(E¼ 2MPa). Applying a magnetic field during curing of

polymer matrix leads to composites, which have a preferred

orientation of the easymagnetisation axis [001] along previous

magnetic field axis. The particles consist of several twin

variants and twin boundaries. The maximal composite strain

can be estimated by the stress-free strain method, yielding the

softer the polymer matrix the higher the magnetic field

induced composite strain. Optical strain measurements on

austenite-polymer-composites show that in a polymer matrix

with a too low Young’s modulus a relocation of particles occur

and the composite bends in magnetic field. Martensi-

te-epoxy-composite, using the stiffer epoxy resin matrix,

shows an MFIS of up to 0.1% after magneto-thermal training,

close to the calculated maximal strain. This is an indication for

a higher mobility of twin boundaries due to magneto-thermal

training. For a further increase of the composite MFIS, the

Young’s modulus of polymer matrix has to be tailored more

precisely and more embedded particles need to be active. The

reason for the decrease in magnetic field induced strain after

each cycle is likely to be a gradual fracture of interfaces

between particles and polymer matrix. This is verified by a

complete failure by adhesive fracture of Ni-Mn-Ga fibres

pulled out of epoxy resin, thereby the apparent shear strength

is in the range of the maximal magnetic force. To overcome

this problem one can prepare composites with epoxysilane

modified fibres. Pull-out tests with epoxysilane treated fibres

show a significant improvement of the interface stability.

Received: April 28, 2011

Final Version: August 26, 2011

Published online: October 18, 2011

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