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On the magnetostructural transition in MnCoGeBx alloy ribbons
Journal of Alloys and Compounds 644 (2015) 1003–1008
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
On the magnetostructural transition in MnCoGeBx alloy ribbons A. Quintana-Nedelcos a,b,⇑, J.L. Sánchez Llamazares a,⇑, H. Flores-Zuñiga a a b
Instituto Potosino de Investigación Científica y Tecnológica (IPICyT) A.C., Camino a la Presa San José No 2055, San Luis Potosí, 78216 S.L.P., Mexico Marmara University, Department of Material and Metalurgy, Kadikoy 3477, Istanbul, Turkey
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Article history: Received 22 January 2015 Received in revised form 6 April 2015 Accepted 1 May 2015 Available online 8 May 2015 Keywords: MnCoGe system alloys Melt-spun ribbons Multiferroic Magneto-structural transition Ferroelastic domain walls Nucleation
a b s t r a c t The magnetostructural transition in the Mn0.96Co1.04GeB0.02 ribbon alloy was investigated. Chemical, structural, microstructural, and magnetic studies were performed on the samples which were annealed at different temperatures. The resulting samples underwent a first-order phase transition in which the characteristic structural transition temperature shows a near-linear and inversely proportional dependence to the annealing temperature. The magnetostructural transition occurs through a simultaneous ferroelastic–magnetic transition between the ferroelastic–paramagnetic and paraelastic–ferromagnetic phases. Our results suggest that the crystal structure of the hexagonal high temperature phase allows the formation of ferroelastic domains, and the domain walls act as the natural nucleation site of the low temperature paraelastic–ferromagnetic phase. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction MnCoGe-based intermetallic alloys transform from a high-temperature Ni2In-type hexagonal (hex) structure to a distorted low-temperature TiNiSi-type orthorhombic (orth) phase [1,2]. Both phases behave as a pure ferromagnetic (FM) material. For the stoichiometric MnCoGe alloy the structural transformation (ST) occurs at T > 398 K [2], far above of the ‘‘temperature window’’ delimited by the Curie temperature of the hexagonal (Thex c ) and orthorhombic (Torth ) phases, 276 K and 355 K respectively [2–5], c in which the magnetic and structural lattices are strongly linked. It has been established that large magnetocaloric effect (MCE) can be obtained due to a coupled magnetostructural first order transformation [6]. In MnCoGe-based alloys, the FM-orth to the paramagnetic (PM)-hex phase transition leads to the giant MCE. Therefore, several efforts have been made to shift the ST of MnCoGe alloys into the Thex to Torth temperature window, including c c (a) the partial atomic substitution of one (or more) of the three main elements of the MnCoGe alloy [4,7–11]; (b) stoichiometry changes [12]; (c) the introduction of elements of small atomic radius into interstitial sites [13]; and (d) the application of physical and chemical pressure [14]. Even though those approaches have been successful in tuning ST, why/how the introduction of lattice ⇑ Corresponding authors at: Marmara University, Department of Material and Metalurgy, Kadikoy 3477, Istanbul, Turkey (A. Quintana-Nedelcos). Tel.: +52 444 2000; fax: +52 444 7269 (J.L. Sánchez Llamazares). E-mail addresses: [email protected] (A. Quintana-Nedelcos), jose.sanchez@ ipicyt.edu.mx (J.L. Sánchez Llamazares). http://dx.doi.org/10.1016/j.jallcom.2015.05.008 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
defects (i.e. atom vacancies, atom substitutions, and deformation of the lattice structure either by the introduction of an interstitial element or by an applied pressure) affect the structural behavior of the MnCoGe system remains unclear. In the present contribution the magnetostructural phase transition of Mn0.96Co1.04GeB0.02 ribbon alloys is study by means of structural, microstructural and magnetic experimental analysis. 2. Experimental A 3 g bulk alloy sample with nominal composition Mn0.96Co1.04GeB0.02 was prepared by arc melting from highly pure starting material (>99.98%). The bulk alloy was melted three times to ensure good homogeneity. As-spun ribbons were obtained by melt spinning under a controlled highly pure argon environment. The copper wheel’s linear speed was 20 ms1. Some amounts of the as-spun ribbons were encapsulated in quartz tubes under an argon atmosphere for a four-hour annealing treatment followed by air-quenching. The annealed temperature (AT) varied from 650 °C to 875 °C. Hereafter, the ribbons annealed at 875 °C, 850 °C, 825 °C, 800 °C, 750 °C, and 650 °C will be referred to as AQ-875, AQ-850, AQ-825, AQ-800, AQ-750, and AQ-650, respectively (series-AQ). The as-quenched ribbons without annealing treatment are named as AQ. Structural, microstructural and, magnetic characterizations were performed for two different sample states: (i) virgin sample: the sample obtained from the air-quenching after the annealing treatment; (ii) cycled sample: samples in which a full DSC cycle was conducted. X-ray powder diffraction (XRD) patterns were obtained with a Bruker AXS model D8 Advance diffractometer using Cu Ka radiation. Phase-transition studies were carried out by differential scanning calorimetry (DSC) using a TA Instruments model Q200. DSC curves for 10 mg samples were measured at a heating/cooling rate of 10 K min1. Microstructural characterization and chemical composition were determined by energy-dispersive X-ray spectroscopy (EDS) using a Phillips model XL-30 scanning electron microscope (SEM). SEM images were taken at room temperature (RT). Magnetization measurements were performed from 10 to 400 K using a Quantum Design PPMS-9T platform with the vibrating sample
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magnetometer module. Field-cooling (FC) and field-heating (FH) thermomagnetic curves were measured for applied static magnetic fields of l0H = 5 mT and l0H = 5 T at a heating or cooling rate of 1 K min1.
3. Results and discussion Fig. 1(a) shows the comparison between the calculated average chemical composition and the nominal composition of all the annealed samples. The chemical composition deviates from the nominal one, and it was calculated as 31.6 at.%, 33.6 at.% and 34.9 at% for Mn, Co and Ge, respectively, giving a global average composition close to Mn0.95CoGe. Fig. 1(b) and (c) show the room-temperature XRD patterns obtained from the as-quenched (AQ) and annealed AQ-750 samples. The AQ-sample was successfully indexed based on a Ni2In-type hexagonal structure. This structure remained in the temperature range from 400 K to 200 K, and was confirmed by the calorimetric study. The XRD pattern showed in Fig. 1(c) corresponds to the low-temperature orthorhombic (orth) phase of the AQ-750 ribbons. The XRD measurement for this sample was performed after cooling down the sample to T = 200 K, before warming it up to room temperature (RT). All the annealed samples transformed from the Ni2In-type hexagonal structure to a distorted low-temperature TiNiSi-type orthorhombic phase, which is in agreement with previous reports [3,10,15]. It has been established that the relationship between the Ni2In-type hexagonal and the TiNiSi-type orthorhombic unit cell pffiffiffi axes is aorth = chex, borth = ahex, and corth = 3ahex [2,16]; which is the only possible misfit angle for the hexagonal to orthorhombic ferroelastic phase transition [17]. Thus, the distortion of the high-symmetry cell occurs primarily in the orthohexagonal ab plane [2,3]. That is, the displacements of atoms must be perpendicular pffiffiffi to the chex axis and along ±bhex (= ± 3ahex ±corth) [18]; this misfit angle favors the hex ? orth ferroelastic transition. Usually, as a consequence of the first-order phase transition, the low-temperature phase or the parent phase (even the untransformed one) allows the formation of a few distinct variants (i.e. the ferroelastic phase) [19]. These variants originate from the spontaneous strain that measures the distortion of the unit cell in contrast to the one in the paraelastic phase [20]. In fact, the ferroelastic domain configuration forms to minimize the macroscopic phase differences between the paraelastic and ferroelastic phases [21]. During the formation of the ferroelastic domain wall between two ferroelastic domains the ideal adjacent domains have to rotate to a certain misfit angle in order to join in a permissible domain wall [17]. Fig. 2(a) and (b) shows the DSC curves for virgin and cycled samples, respectively. However, it is important to note that the typical structural transition temperature is inversely proportional to the annealing temperature. Fig. 2(c) shows the near-linear dependence of the starting and finishing temperatures for the direct and reverse martensitic transformation (Ms, Mf, Af, and As respectively) with the treatment temperature. These values were taken from the second DSC cycle measurement. For a variety of shape memory alloys it has been established that the martensitic starting temperature depends on the average valence electron concentration per atom (e/a) [12,22] and on the average grain size (hdi) [23,24]. However, our samples exhibit a high degree of homogeneity in their average elemental composition and microstructure. Hence, the Ms dependence on both e/a and hdi can be neglected. Additionally, Ms can be further tuned through modification of the annealing process. In this sense, it is established that depending on the annealing temperature and time the stress that is formed during the quenching process can be relaxed to some extent, in which the atom site/ordering and lattice symmetry can show a modification [25–27]. A decrease in Ms caused by an increase of the ageing temperature conducted in Mn1xCrxCoGe alloys was explained by the
strong effect that the annealing temperature had on the order degree of the high temperature phase [28]. This interpretation was made in accordance with the observation of similar results in different shape memory alloys undergoing a martensitic transformation. However, the key-role played by the impurities in the MnCoGe based system alloys was not considered. In this sense, a quick explanation may be attempted in order to find a direct correlation between the lattice defects concentration and Ms. Moreover, it is noteworthy that all ribbons of our study were fabricated from a single bulk master alloy. In consequence, even though a small variation in the lattice defects concentration (i.e. boron atoms and vacancies) may exist among the ribbons, this is expected to be ageing independent. On the other hand, a high density of dislocations and laminar microstructure along with some preferential surface planes were observed in Mn1xCrxCoGe alloys by transmission electron microscopy (TEM) [28]. Thus, the inhomogeneous distribution of the lattice defects seems to be the key-factor for understanding the Ms shift. In such a sense, the dislocation distribution can be highly affected by the annealing treatment. The latter can be separated into a two-steps process. (Annealing) The high annealing temperatures allow a bigger and faster dislocation mobility, as well as lower stress and a more homogeneous distribution of the stress among the lattice. (Cooling) The slow cooling from the annealing temperature permits defects diffusion from their high annealed temperature random position to the more stable position along the preferential surface planes. With the decrease of the cooling rate the diffusion of the dislocations is favored. If the alloy is cooled rapidly (i.e. water quenching) the point defects do not have the time to diffuse and are held in their high temperature lattice positions, as a consequence no martensitic transformation is observed [5,28]. In our case all samples were air-quenched, however, when the cooling periods are delayed they have the same effect as slower cooling rates (the rising of the annealing temperature increases the cooling time). In conclusion, as the lattice defects density increases along the preferential orthohexagonal ab planes (with the increases of the annealing temperature), the hex ? orth transition temperature decreases. Ms and Mf of samples AQ-875, AQ-850, AQ-825, AQ-800, and AQ-750 present a significant shift (66 K) to higher temperatures when they are compared with those found for first and second DSC-curves. There were not observed any significant variations on the transition temperatures in further cycles (i.e. second, third and fourth cycles show same futures). The Ms and Mf shift was not observed in AQ-650 ribbons. Formally, this result can be interpreted as follows: the AQ-650 sample suffers the transformation before it reaches the RT during the cooling step after annealing (having the same effect of the first DSC-cycle for the rest of the samples). As a result, the sample needs to be structurally transformed once in order to reach more favorable conditions for further transformations. In which probably, the dislocations relocate toward the transformation preferential planes. The thermal hysteresis of cycled samples (determined from the second DSC cycle) was found to be > 24 K. However, even if the considerable thermal hysteresis is not suitable for practical applications, this is an extrinsic property, can be reduced by processing [29,30]. Figs. 3(a) and (b) show the DSC and low-field magnetization curves of the AQ-750 virgin and cycled sample, respectively. Figs. 3(c) and (d) show the DSC and low-field magnetization curves of the AQ-850 virgin and cycled sample, respectively. As shown, the shift in the cooling peak of the structural transition is fixed with a shift in the field-cooling curve of the magnetic transformation. The latter is due to the coupled magneto-structural behavior of the transformation. The thermal hysteresis of the magnetization curves also indicates the first-order behavior of the transition. The Curie temperature of the hexagonal phase was derived from the
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Fig. 1. Figure (a) shows the average atomic concentration and the standard deviation calculated from EDS for Co, Mn, and Ge elements. Element atomic concentration is plotted as a function of the treatment temperature. The dashed lines represent the nominal elemental atomic concentration (i.e. 34.67 at.% Co, 32 at.% Mn and 33.33 at.% Ge). Figure (b) and (c) shows the room temperature XRD patterns for AQ (b) and AQ-750 (c) alloy ribbons.
Fig. 2. Figures (a) and (b) show the DSC curves for cooling and heating for the first and second cycles of measurement, respectively. Figure (c) shows the austenitic and martensitic starting and finishing transition temperature dependence with the annealed treatment temperature for treated ribbons. Data point was taken from figure (b). Dashed lines represent the Curie temperature of the orthorhombic and hexagonal phases for the stoichiometric MnCoGe alloy. Figure (d) and (e) correspond to the first and second cycle of the DSC curves of the ribbons (d) and manually mill (e) AQ-800 samples.
second cycle magnetization curve. The given result for samples AQ-750 and AQ-850 are Thex 258 K and 262 K, respectively, c which are similar to those reported for close-stoichiometric bulk alloys Mn0.95CoGe (265 K) [2], and Mn1xCoGe with x = 0.045, and x = 0.050 (261 K, and 260 K respectively) [12]. Note that for the virgin AQ-850 sample (Fig. 3(c)), the structural transition on cooling occurs below the Curie temperature of the high temperature hexagonal phase. In this regard, when the sample is cooled down from a high temperature, it first transforms magnetically from a PM-hex state to a FM-hex phase, then when the temperature reaches the structural transition temperature, it structurally transforms from the FM-hex phase to the FM-orth phase. In the heating pathway of the first cycle the sample transforms from the FM-orth phase directly to the PM-hex phase. When further
thermal cycles take place (i.e. a second thermal cycle, Fig. 3(d)) there is a transformation, either on the cooling or heating process, between the FM-orth and the PM-hex phases. Fig. 4 shows the ribbon surface micrographs of the samples taken under different conditions. SEM images of AQ-650, AQ-750, AQ-800, AQ-825, AQ-850, and AQ-875 virgin samples are shown from Figs. 4(a) to (f). Average grain size of 5 ± 3 lm (calculated as the average grain diameter) proves to be independent of the annealing temperature. Among the virgin samples, the AQ-650 is the only one which shows fractures on its surface, as it was the only sample transformed into the orthorhombic structure phase when cooling from the annealing temperature. The reason for severe cracking in the sample is the large negative volume change (4%) which accompanies the hex ? orth transition in
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Fig. 3. First (a) and second (b) cycle of the magnetization and DSC curves for samples AQ-750 and first (c) and second (d) cycle of the magnetization and DSC curves for samples AQ-850. The Curie temperature of the hexagonal phase is pointed out by dashed arrows on the M(T) curves.
Fig. 4. Typical SEM images of the ribbons surface microstructure. Images from (a) to (f) were taken on virgin samples. Images (g), (h), and (i) were taken on cycled samples.
MnCoGe-based alloys [2]. In order to remove the variable introduced by the cracking of the material during the transformation in the shift of the characteristic transition temperatures some amounts of virgin ribbons were milled manually and then subjected to the same calorimetric analysis applied to the ribbons. Fig. 2(d) and (e) shows the first and second DSC cycles of the ribbons and milled AQ-800 virgin alloys, respectively. In fact, there were no differences observed in the characteristics transition temperatures between both ribbon and milled sample. It should be noted, that there was a difference in the first hex ? orth
transformation of the milled sample. The inherent small peaks associated to the fracture process are not showed. The straight lines observed inside the grains in the untransformed virgin samples (Fig. 4(b)–(f)) have been described as ferroelastic domain walls (FEDW) [2] in accordance with the ferroelastic transition type between the hexagonal P63/mmc and the orthorhombic Pnma space group [31]. From the absence of those lines in the virgin sample AQ-650 we can say that the appearance of such FEDWs occur in the high temperature hexagonal structure. Also, a monotone increase of the FEDW density with the decrease
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of the annealing temperature is observed in untransformed virgin samples. This observation can be interpreted as follow: the FEDW density increases when the temperature draws near the structural transition. In fact, this result is in complete agreement with the findings of Wenyuan et al., who proved that the number of domains becomes larger as the temperature approaches the fer roelastic–paraelastic transition temperature [32]. Fig. 4(g)–(i) show the room temperature micrographs of AQ-800, AQ-825, and AQ-850 cycled samples, respectively. SEM images of AQ-800, AQ-825, and AQ-850 cycled alloys correspond to the orth, orth + hex and, hex structural phases, respectively. The absence of the FEDWs in the orthorhombic structure of the AQ-800 cycled sample (Fig. 4(g)) confirms the previous statement that identified the high-temperature hexagonal structure and the low-temperature orthorhombic structure as the ferroelastic and paraelastic phases, respectively. Fig. 4(h) shows the co-existence of the ferroelastic and paraelastic phases as the transition proceeds. And, Fig. 4(i) shows that the ferroelastic variants develop when the sample transforms from the orth to the hex phase. Fig. 5 schematically represents the process of nucleation and growth of the paraelastic-orth phase into the ferroelastic-hex matrix. Regarding to this, it is known that the mobility of the ferroelastic domain walls is determined by its thickness [33]. Thin domain walls interact with the point defects in the crystal structure to a greater degree than thicker domain walls, therefore thicker domain walls are more mobile [34]. However, ferroelasticity is often accompanied by other phenomena such as the change in the magnetic order [35]. Therefore if a ferroelastic material is also ferromagnetic the structure of the domain walls will be a consequence of gradients in the magnetic order and strain simultaneously [33]. The FEDW has a characteristic thickness of the order of the cell parameter [36], while the walls of the magnetic domains are much wider, in the order of nanometers [34]. Thus, the FEDW thickness increases between the limits given by the magnetic and elastic cases pure domains [33]. In annealed ribbons, when an elasto-magnetic transition is to be considered, it is expected that the FEDWs have a relatively high thickness which promote the necessary FEDW mobility to ensure the paraelastic-orth domain growth. Additionally, the ferroelastic domains form in order to reduce the internal elastic energy [21,37] and its size responds to an equilibrium between the domain energy and the domain-wall energy [21]. The gradient stress necessary to form the domains can be originated from the self-stress imposed, for example, from the intrinsic surface tension [21], and from material inhomogeneities such as lattice defects and polycrystalline structure [37]. This could be the explanation of why the different approaches given in the introduction have all been successful at tuning the structural transformation into the temperature window; all of them introduce the necessary lattice defects to originate the gradient stress which promotes the elastic domain formation. 4. Conclusions In conclusion, our samples can be considered as a multiferroic material which undergoes a reversible first-order phase transition between the paraelastic–ferromagnetic (orth) and the ferroelastic– paramagnetic (hex) phases. The discussed results pointed out to the FEDWs as the natural orthorhombic nucleation sites into the hexagonal matrix in which a single FEDW divides two regions characterized by a different orientation. That is, FEDWs act as the low-temperature nucleation site and the new phase grows as the domain wall propagates through the crystal. Our results could lead to an important advantage from the potential application standpoint. The engineering design of the material with some required issues can be made on the basis of the knowledge of the extrinsic
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Fig. 5. Characteristic on cooling pathway DSC curve upon which schematically represents the formation and growth of domain boundaries during the ferroelastic (hex) to paraelastic (orth) transition as the temperature decreases.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
On the magnetostructural transition in MnCoGeBx alloy ribbons A. Quintana-Nedelcos a,b,⇑, J.L. Sánchez Llamazares a,⇑, H. Flores-Zuñiga a a b
Instituto Potosino de Investigación Científica y Tecnológica (IPICyT) A.C., Camino a la Presa San José No 2055, San Luis Potosí, 78216 S.L.P., Mexico Marmara University, Department of Material and Metalurgy, Kadikoy 3477, Istanbul, Turkey
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Article history: Received 22 January 2015 Received in revised form 6 April 2015 Accepted 1 May 2015 Available online 8 May 2015 Keywords: MnCoGe system alloys Melt-spun ribbons Multiferroic Magneto-structural transition Ferroelastic domain walls Nucleation
a b s t r a c t The magnetostructural transition in the Mn0.96Co1.04GeB0.02 ribbon alloy was investigated. Chemical, structural, microstructural, and magnetic studies were performed on the samples which were annealed at different temperatures. The resulting samples underwent a first-order phase transition in which the characteristic structural transition temperature shows a near-linear and inversely proportional dependence to the annealing temperature. The magnetostructural transition occurs through a simultaneous ferroelastic–magnetic transition between the ferroelastic–paramagnetic and paraelastic–ferromagnetic phases. Our results suggest that the crystal structure of the hexagonal high temperature phase allows the formation of ferroelastic domains, and the domain walls act as the natural nucleation site of the low temperature paraelastic–ferromagnetic phase. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction MnCoGe-based intermetallic alloys transform from a high-temperature Ni2In-type hexagonal (hex) structure to a distorted low-temperature TiNiSi-type orthorhombic (orth) phase [1,2]. Both phases behave as a pure ferromagnetic (FM) material. For the stoichiometric MnCoGe alloy the structural transformation (ST) occurs at T > 398 K [2], far above of the ‘‘temperature window’’ delimited by the Curie temperature of the hexagonal (Thex c ) and orthorhombic (Torth ) phases, 276 K and 355 K respectively [2–5], c in which the magnetic and structural lattices are strongly linked. It has been established that large magnetocaloric effect (MCE) can be obtained due to a coupled magnetostructural first order transformation [6]. In MnCoGe-based alloys, the FM-orth to the paramagnetic (PM)-hex phase transition leads to the giant MCE. Therefore, several efforts have been made to shift the ST of MnCoGe alloys into the Thex to Torth temperature window, including c c (a) the partial atomic substitution of one (or more) of the three main elements of the MnCoGe alloy [4,7–11]; (b) stoichiometry changes [12]; (c) the introduction of elements of small atomic radius into interstitial sites [13]; and (d) the application of physical and chemical pressure [14]. Even though those approaches have been successful in tuning ST, why/how the introduction of lattice ⇑ Corresponding authors at: Marmara University, Department of Material and Metalurgy, Kadikoy 3477, Istanbul, Turkey (A. Quintana-Nedelcos). Tel.: +52 444 2000; fax: +52 444 7269 (J.L. Sánchez Llamazares). E-mail addresses: [email protected] (A. Quintana-Nedelcos), jose.sanchez@ ipicyt.edu.mx (J.L. Sánchez Llamazares). http://dx.doi.org/10.1016/j.jallcom.2015.05.008 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
defects (i.e. atom vacancies, atom substitutions, and deformation of the lattice structure either by the introduction of an interstitial element or by an applied pressure) affect the structural behavior of the MnCoGe system remains unclear. In the present contribution the magnetostructural phase transition of Mn0.96Co1.04GeB0.02 ribbon alloys is study by means of structural, microstructural and magnetic experimental analysis. 2. Experimental A 3 g bulk alloy sample with nominal composition Mn0.96Co1.04GeB0.02 was prepared by arc melting from highly pure starting material (>99.98%). The bulk alloy was melted three times to ensure good homogeneity. As-spun ribbons were obtained by melt spinning under a controlled highly pure argon environment. The copper wheel’s linear speed was 20 ms1. Some amounts of the as-spun ribbons were encapsulated in quartz tubes under an argon atmosphere for a four-hour annealing treatment followed by air-quenching. The annealed temperature (AT) varied from 650 °C to 875 °C. Hereafter, the ribbons annealed at 875 °C, 850 °C, 825 °C, 800 °C, 750 °C, and 650 °C will be referred to as AQ-875, AQ-850, AQ-825, AQ-800, AQ-750, and AQ-650, respectively (series-AQ). The as-quenched ribbons without annealing treatment are named as AQ. Structural, microstructural and, magnetic characterizations were performed for two different sample states: (i) virgin sample: the sample obtained from the air-quenching after the annealing treatment; (ii) cycled sample: samples in which a full DSC cycle was conducted. X-ray powder diffraction (XRD) patterns were obtained with a Bruker AXS model D8 Advance diffractometer using Cu Ka radiation. Phase-transition studies were carried out by differential scanning calorimetry (DSC) using a TA Instruments model Q200. DSC curves for 10 mg samples were measured at a heating/cooling rate of 10 K min1. Microstructural characterization and chemical composition were determined by energy-dispersive X-ray spectroscopy (EDS) using a Phillips model XL-30 scanning electron microscope (SEM). SEM images were taken at room temperature (RT). Magnetization measurements were performed from 10 to 400 K using a Quantum Design PPMS-9T platform with the vibrating sample
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magnetometer module. Field-cooling (FC) and field-heating (FH) thermomagnetic curves were measured for applied static magnetic fields of l0H = 5 mT and l0H = 5 T at a heating or cooling rate of 1 K min1.
3. Results and discussion Fig. 1(a) shows the comparison between the calculated average chemical composition and the nominal composition of all the annealed samples. The chemical composition deviates from the nominal one, and it was calculated as 31.6 at.%, 33.6 at.% and 34.9 at% for Mn, Co and Ge, respectively, giving a global average composition close to Mn0.95CoGe. Fig. 1(b) and (c) show the room-temperature XRD patterns obtained from the as-quenched (AQ) and annealed AQ-750 samples. The AQ-sample was successfully indexed based on a Ni2In-type hexagonal structure. This structure remained in the temperature range from 400 K to 200 K, and was confirmed by the calorimetric study. The XRD pattern showed in Fig. 1(c) corresponds to the low-temperature orthorhombic (orth) phase of the AQ-750 ribbons. The XRD measurement for this sample was performed after cooling down the sample to T = 200 K, before warming it up to room temperature (RT). All the annealed samples transformed from the Ni2In-type hexagonal structure to a distorted low-temperature TiNiSi-type orthorhombic phase, which is in agreement with previous reports [3,10,15]. It has been established that the relationship between the Ni2In-type hexagonal and the TiNiSi-type orthorhombic unit cell pffiffiffi axes is aorth = chex, borth = ahex, and corth = 3ahex [2,16]; which is the only possible misfit angle for the hexagonal to orthorhombic ferroelastic phase transition [17]. Thus, the distortion of the high-symmetry cell occurs primarily in the orthohexagonal ab plane [2,3]. That is, the displacements of atoms must be perpendicular pffiffiffi to the chex axis and along ±bhex (= ± 3ahex ±corth) [18]; this misfit angle favors the hex ? orth ferroelastic transition. Usually, as a consequence of the first-order phase transition, the low-temperature phase or the parent phase (even the untransformed one) allows the formation of a few distinct variants (i.e. the ferroelastic phase) [19]. These variants originate from the spontaneous strain that measures the distortion of the unit cell in contrast to the one in the paraelastic phase [20]. In fact, the ferroelastic domain configuration forms to minimize the macroscopic phase differences between the paraelastic and ferroelastic phases [21]. During the formation of the ferroelastic domain wall between two ferroelastic domains the ideal adjacent domains have to rotate to a certain misfit angle in order to join in a permissible domain wall [17]. Fig. 2(a) and (b) shows the DSC curves for virgin and cycled samples, respectively. However, it is important to note that the typical structural transition temperature is inversely proportional to the annealing temperature. Fig. 2(c) shows the near-linear dependence of the starting and finishing temperatures for the direct and reverse martensitic transformation (Ms, Mf, Af, and As respectively) with the treatment temperature. These values were taken from the second DSC cycle measurement. For a variety of shape memory alloys it has been established that the martensitic starting temperature depends on the average valence electron concentration per atom (e/a) [12,22] and on the average grain size (hdi) [23,24]. However, our samples exhibit a high degree of homogeneity in their average elemental composition and microstructure. Hence, the Ms dependence on both e/a and hdi can be neglected. Additionally, Ms can be further tuned through modification of the annealing process. In this sense, it is established that depending on the annealing temperature and time the stress that is formed during the quenching process can be relaxed to some extent, in which the atom site/ordering and lattice symmetry can show a modification [25–27]. A decrease in Ms caused by an increase of the ageing temperature conducted in Mn1xCrxCoGe alloys was explained by the
strong effect that the annealing temperature had on the order degree of the high temperature phase [28]. This interpretation was made in accordance with the observation of similar results in different shape memory alloys undergoing a martensitic transformation. However, the key-role played by the impurities in the MnCoGe based system alloys was not considered. In this sense, a quick explanation may be attempted in order to find a direct correlation between the lattice defects concentration and Ms. Moreover, it is noteworthy that all ribbons of our study were fabricated from a single bulk master alloy. In consequence, even though a small variation in the lattice defects concentration (i.e. boron atoms and vacancies) may exist among the ribbons, this is expected to be ageing independent. On the other hand, a high density of dislocations and laminar microstructure along with some preferential surface planes were observed in Mn1xCrxCoGe alloys by transmission electron microscopy (TEM) [28]. Thus, the inhomogeneous distribution of the lattice defects seems to be the key-factor for understanding the Ms shift. In such a sense, the dislocation distribution can be highly affected by the annealing treatment. The latter can be separated into a two-steps process. (Annealing) The high annealing temperatures allow a bigger and faster dislocation mobility, as well as lower stress and a more homogeneous distribution of the stress among the lattice. (Cooling) The slow cooling from the annealing temperature permits defects diffusion from their high annealed temperature random position to the more stable position along the preferential surface planes. With the decrease of the cooling rate the diffusion of the dislocations is favored. If the alloy is cooled rapidly (i.e. water quenching) the point defects do not have the time to diffuse and are held in their high temperature lattice positions, as a consequence no martensitic transformation is observed [5,28]. In our case all samples were air-quenched, however, when the cooling periods are delayed they have the same effect as slower cooling rates (the rising of the annealing temperature increases the cooling time). In conclusion, as the lattice defects density increases along the preferential orthohexagonal ab planes (with the increases of the annealing temperature), the hex ? orth transition temperature decreases. Ms and Mf of samples AQ-875, AQ-850, AQ-825, AQ-800, and AQ-750 present a significant shift (66 K) to higher temperatures when they are compared with those found for first and second DSC-curves. There were not observed any significant variations on the transition temperatures in further cycles (i.e. second, third and fourth cycles show same futures). The Ms and Mf shift was not observed in AQ-650 ribbons. Formally, this result can be interpreted as follows: the AQ-650 sample suffers the transformation before it reaches the RT during the cooling step after annealing (having the same effect of the first DSC-cycle for the rest of the samples). As a result, the sample needs to be structurally transformed once in order to reach more favorable conditions for further transformations. In which probably, the dislocations relocate toward the transformation preferential planes. The thermal hysteresis of cycled samples (determined from the second DSC cycle) was found to be > 24 K. However, even if the considerable thermal hysteresis is not suitable for practical applications, this is an extrinsic property, can be reduced by processing [29,30]. Figs. 3(a) and (b) show the DSC and low-field magnetization curves of the AQ-750 virgin and cycled sample, respectively. Figs. 3(c) and (d) show the DSC and low-field magnetization curves of the AQ-850 virgin and cycled sample, respectively. As shown, the shift in the cooling peak of the structural transition is fixed with a shift in the field-cooling curve of the magnetic transformation. The latter is due to the coupled magneto-structural behavior of the transformation. The thermal hysteresis of the magnetization curves also indicates the first-order behavior of the transition. The Curie temperature of the hexagonal phase was derived from the
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Fig. 1. Figure (a) shows the average atomic concentration and the standard deviation calculated from EDS for Co, Mn, and Ge elements. Element atomic concentration is plotted as a function of the treatment temperature. The dashed lines represent the nominal elemental atomic concentration (i.e. 34.67 at.% Co, 32 at.% Mn and 33.33 at.% Ge). Figure (b) and (c) shows the room temperature XRD patterns for AQ (b) and AQ-750 (c) alloy ribbons.
Fig. 2. Figures (a) and (b) show the DSC curves for cooling and heating for the first and second cycles of measurement, respectively. Figure (c) shows the austenitic and martensitic starting and finishing transition temperature dependence with the annealed treatment temperature for treated ribbons. Data point was taken from figure (b). Dashed lines represent the Curie temperature of the orthorhombic and hexagonal phases for the stoichiometric MnCoGe alloy. Figure (d) and (e) correspond to the first and second cycle of the DSC curves of the ribbons (d) and manually mill (e) AQ-800 samples.
second cycle magnetization curve. The given result for samples AQ-750 and AQ-850 are Thex 258 K and 262 K, respectively, c which are similar to those reported for close-stoichiometric bulk alloys Mn0.95CoGe (265 K) [2], and Mn1xCoGe with x = 0.045, and x = 0.050 (261 K, and 260 K respectively) [12]. Note that for the virgin AQ-850 sample (Fig. 3(c)), the structural transition on cooling occurs below the Curie temperature of the high temperature hexagonal phase. In this regard, when the sample is cooled down from a high temperature, it first transforms magnetically from a PM-hex state to a FM-hex phase, then when the temperature reaches the structural transition temperature, it structurally transforms from the FM-hex phase to the FM-orth phase. In the heating pathway of the first cycle the sample transforms from the FM-orth phase directly to the PM-hex phase. When further
thermal cycles take place (i.e. a second thermal cycle, Fig. 3(d)) there is a transformation, either on the cooling or heating process, between the FM-orth and the PM-hex phases. Fig. 4 shows the ribbon surface micrographs of the samples taken under different conditions. SEM images of AQ-650, AQ-750, AQ-800, AQ-825, AQ-850, and AQ-875 virgin samples are shown from Figs. 4(a) to (f). Average grain size of 5 ± 3 lm (calculated as the average grain diameter) proves to be independent of the annealing temperature. Among the virgin samples, the AQ-650 is the only one which shows fractures on its surface, as it was the only sample transformed into the orthorhombic structure phase when cooling from the annealing temperature. The reason for severe cracking in the sample is the large negative volume change (4%) which accompanies the hex ? orth transition in
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Fig. 3. First (a) and second (b) cycle of the magnetization and DSC curves for samples AQ-750 and first (c) and second (d) cycle of the magnetization and DSC curves for samples AQ-850. The Curie temperature of the hexagonal phase is pointed out by dashed arrows on the M(T) curves.
Fig. 4. Typical SEM images of the ribbons surface microstructure. Images from (a) to (f) were taken on virgin samples. Images (g), (h), and (i) were taken on cycled samples.
MnCoGe-based alloys [2]. In order to remove the variable introduced by the cracking of the material during the transformation in the shift of the characteristic transition temperatures some amounts of virgin ribbons were milled manually and then subjected to the same calorimetric analysis applied to the ribbons. Fig. 2(d) and (e) shows the first and second DSC cycles of the ribbons and milled AQ-800 virgin alloys, respectively. In fact, there were no differences observed in the characteristics transition temperatures between both ribbon and milled sample. It should be noted, that there was a difference in the first hex ? orth
transformation of the milled sample. The inherent small peaks associated to the fracture process are not showed. The straight lines observed inside the grains in the untransformed virgin samples (Fig. 4(b)–(f)) have been described as ferroelastic domain walls (FEDW) [2] in accordance with the ferroelastic transition type between the hexagonal P63/mmc and the orthorhombic Pnma space group [31]. From the absence of those lines in the virgin sample AQ-650 we can say that the appearance of such FEDWs occur in the high temperature hexagonal structure. Also, a monotone increase of the FEDW density with the decrease
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of the annealing temperature is observed in untransformed virgin samples. This observation can be interpreted as follow: the FEDW density increases when the temperature draws near the structural transition. In fact, this result is in complete agreement with the findings of Wenyuan et al., who proved that the number of domains becomes larger as the temperature approaches the fer roelastic–paraelastic transition temperature [32]. Fig. 4(g)–(i) show the room temperature micrographs of AQ-800, AQ-825, and AQ-850 cycled samples, respectively. SEM images of AQ-800, AQ-825, and AQ-850 cycled alloys correspond to the orth, orth + hex and, hex structural phases, respectively. The absence of the FEDWs in the orthorhombic structure of the AQ-800 cycled sample (Fig. 4(g)) confirms the previous statement that identified the high-temperature hexagonal structure and the low-temperature orthorhombic structure as the ferroelastic and paraelastic phases, respectively. Fig. 4(h) shows the co-existence of the ferroelastic and paraelastic phases as the transition proceeds. And, Fig. 4(i) shows that the ferroelastic variants develop when the sample transforms from the orth to the hex phase. Fig. 5 schematically represents the process of nucleation and growth of the paraelastic-orth phase into the ferroelastic-hex matrix. Regarding to this, it is known that the mobility of the ferroelastic domain walls is determined by its thickness [33]. Thin domain walls interact with the point defects in the crystal structure to a greater degree than thicker domain walls, therefore thicker domain walls are more mobile [34]. However, ferroelasticity is often accompanied by other phenomena such as the change in the magnetic order [35]. Therefore if a ferroelastic material is also ferromagnetic the structure of the domain walls will be a consequence of gradients in the magnetic order and strain simultaneously [33]. The FEDW has a characteristic thickness of the order of the cell parameter [36], while the walls of the magnetic domains are much wider, in the order of nanometers [34]. Thus, the FEDW thickness increases between the limits given by the magnetic and elastic cases pure domains [33]. In annealed ribbons, when an elasto-magnetic transition is to be considered, it is expected that the FEDWs have a relatively high thickness which promote the necessary FEDW mobility to ensure the paraelastic-orth domain growth. Additionally, the ferroelastic domains form in order to reduce the internal elastic energy [21,37] and its size responds to an equilibrium between the domain energy and the domain-wall energy [21]. The gradient stress necessary to form the domains can be originated from the self-stress imposed, for example, from the intrinsic surface tension [21], and from material inhomogeneities such as lattice defects and polycrystalline structure [37]. This could be the explanation of why the different approaches given in the introduction have all been successful at tuning the structural transformation into the temperature window; all of them introduce the necessary lattice defects to originate the gradient stress which promotes the elastic domain formation. 4. Conclusions In conclusion, our samples can be considered as a multiferroic material which undergoes a reversible first-order phase transition between the paraelastic–ferromagnetic (orth) and the ferroelastic– paramagnetic (hex) phases. The discussed results pointed out to the FEDWs as the natural orthorhombic nucleation sites into the hexagonal matrix in which a single FEDW divides two regions characterized by a different orientation. That is, FEDWs act as the low-temperature nucleation site and the new phase grows as the domain wall propagates through the crystal. Our results could lead to an important advantage from the potential application standpoint. The engineering design of the material with some required issues can be made on the basis of the knowledge of the extrinsic
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Fig. 5. Characteristic on cooling pathway DSC curve upon which schematically represents the formation and growth of domain boundaries during the ferroelastic (hex) to paraelastic (orth) transition as the temperature decreases.
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