Journal of Alloys and Compounds, 205 (1994) 157-163 157 JALCOM 969

Thermomechanical treatment and heat treatment of Nd2.sFe13CoCl.a25Bo.o75 bulk material

Ch. H e l l w i g lnstitut fiir Kristallographie und Petrographie, ETHZ, CH-8092 Ziirich (Switzerland)

(Received September 5, 1993)

Abstract

Nd2.sFe13CoCl.4~Bo.o7 5 is considered as new permanent magnet material of technical interest. Time-temperature- transition diagrams were investigated for 800, 850, 900 and 950 °C. The starting microstructure consists mainly of dendrites of free iron with an intergranular Nd-rich phase. The desired Nd2Fe14B-type phase (q~ phase, tetragonal) forms at the interfaces of iron/Nd-rich. After 15 rain at 850 °C more than 50% of the material consists of the • phase. The application of a thermomechanical treatment (hot deforming) causes the q) phase crystallites to break and to align in the flow direction of the material. A large amount of Nd-rich phase can be pressed out of the sample during deformation. Further experiments are necessary to prevent this. Another technique to obtain anisotropic bulk material is the heat treatment in a temperature gradient.

1. Introduction

A promising group of permanent magnet materials, R2FelaB (R is a rare earth metal), was discovered in 1984 [1]. Magnetic and structural properties of these ternary compounds are reviewed in ref. 2. Substituting boron by carbon led to another promising group of permanent magnet materials, the isotypic compounds R2Fe14C [3-6]. Their Curie temperatures are slightly lower than those of the corresponding R2Fe~4B systems (e.g. 535 K for Nd2Fe~4C [7], compared with 585 K for Nd2Fe~4B [2]). However, the R2Fe14C compounds are formed peritectically and undergo a solid-state trans- formation. At higher temperatures the A phase (rhom- bohedral space group R3m, isotypic to Th2Zn17) is stable, while at lower temperatures the • phase (te- tragonal space group P42/mnm, isotypic to Nd2Fe14B) is stable. This offers new possibilities for obtaining high coercivity bulk material without the necessity of em- ploying the powder metallurgical route. A lot of in- vestigations were made on hot deforming of Nd-Fe-B or Pr-Fe-B alloys [8-13]. Furthermore the Nd-Fe-C system was investigated by different research groups [14-17].

The aim of the present work is to investigate the time-temperature-transition (ttt) behaviour of Nd2.: Fe13CoCIA25B0.o75 and to investigate the possibilities of applying a thermomechanical treatment or a special heat treatment in order to obtain anisotropic bulk material.

An orientation of the crystallites during the appli- cation of heat and pressure could be observed in Nd2FeI4B compounds [18]. An orientation of crystallites during extrudation was also observed in Nd-rich Nd2Fe14B [19].

2. Experimental details

The alloys were prepared by arc melting from starting materials of at least 99.9% purity. We started with an Nd excess of 3% to equalize evaporation losses during arc melting. The composition Nd2.sFelafoCl.425Bo.o75 was chosen for preparation because it passes a nearly complete phase transition between 1080 °C and 850 °C [20]. The samples were wrapped into tantalum foil, sealed into an evacuated quartz tube and annealed at 1150 °C for 15 min to obtain a well-defined starting microstructure. The samples were subsequently an- nealed at different temperatures for different times without cooling after the heat treatment at 1150 °C. For annealing temperatures and times see Table 1. The microstructures of the annealed samples were studied by standard metallographic techniques. The samples were etched in an acetic acid/ethanol solution for phase distinction. Additional microhardness mea- surements were made for phase identification.

X-ray powder photographs were obtained using a Guinier focusing camera (Jagodzinski type, Fe-Ka ra- diation) with Si as an internal standard. The intensities

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158 Ch. Hellwig / Thermomechanical treatment of Ndz~Fe13CoCi.425Bo.oz5

TABLE 1. Summary of the t ime-temperature-transit ion inves- tigations of Nd2.sFe13CoCI.42sBo.075. The standard deviation is about 5% of the value

Temperature Time a-iron Nd-rich q~ phase A phase (°C) (min) (%) phase (%) (%)

(%)

800 2 47 23 0.5 30 800 5 39 19 21 21 800 15 40 19 25 16 800 30 23 10 36 32 800 70 19 11 39 32 800 300 10 12 62 16 800 1440 3 12 71 15

850 1 53 29 0.5 18 850 2 37 25 13 25 850 5 41 22 25 12 850 9 29 17 46 8 850 15 38 19 37 6 850 30 25 12 50 13 850 60 22 8 55 15 850 300 11 7 62 20 850 1440 3 6 73 18

900 2 50 31 10 10 900 3.5 37 22 19 22 900 6 30 16 36 18 900 10 35 16 40 9 900 60 15 13 52 20 900 300 15 9 49 27 900 1440 15 10 38 37

950 2 56 21 6 17 950 10 31 19 37 13 950 60 19 11 41 29 950 300 19 11 34 36 950 1440 15 8 34 43

To investigate for preferred orientation the bulk samples were measured in a powder diffractometer (Cu-Ka radiation). Only crystal planes parallel to the surface contribute to X-ray diffraction due to the Bragg-Brentano geometry. The angles between 35 ° and 46 ° 20 were observed.

The saturation magnetization was measured with a moving sample magnetometer at 4.2 K with B up to 5 T. About 80 mg bulk material with cubic shape was used for these measurements. Values for the saturation magnetization Ms were derived from the higher-field part of the M(B) curves, by extrapolation to zero field. Samples that show a hysteresis were measured parallel and perpendicular to the flow direction of the material during hot deforming. For the other samples no dif- ferences in Ms were expected.

3. Results

3.1. Time-temperature-transition behaviour The results of the ttt investigations are listed in Table

1 and four diagrams are shown in Figs. 1-4. The irregularities (especially the iron content at 850 °C) are due to lower accuracy in measurement or to in- homogeneity of the samples. Nevertheless the tendency of the ttt behaviour is quite obvious. Two micrographs show the formation of the intermetallic phases at 850 °C (Figs. 5 and 6). The • phase and the A phase coexist. The higher the temperature, the more h phase is formed. A temperature of 850 °C was chosen for further investigations.

were measured by means of an automatic Guinier-film scanner. Intensities of two reflections of each phase (i.e. the q~ phase and the A phase) were used to calculate the phase content using a method described by Hill [21]. Only the ratio qb:A was determined, not the absolute content. Comparisons between these measurements and standard metallographic technique measurements yield a very good agreement. The amounts of a-iron and Nd-rich phase were not measured quantitatively by X- ray diffraction because of preparation reasons. There- fore these amounts had to be determined by standard metallographic techniques that yield less accurate re- sults.

Three techniques were applied in order to influence the crystallization of the q~ phase: hot deforming during recrystallization between two parallel plates at 850 °C with a maximum of 60 kN and at around 1 mm s-1; annealing in a temperature gradient of 22 °C mm-1 in a zone melting apparatus and annealing in an in- duction furnace at a frequency of 400 kHz (i.e. in a high frequency magnetic field).

3.2. Hot deformation during crystallization The samples were deformed as listed in Table 2.

Afterwards the samples were characterized by micro- graphs and X-ray measurements, annealed at 850 °C for 3 weeks and characterized again. The deformed

':i 8oooq z - f ""'r°n :~ 80

40 "" O - p h a s e 30

-8> 2010 ~ Nd-rich~ ~ x'--T~ ~ 0 phose ± ~

1 10 100 100( annea l ing t ime [ ra in ]

Fig. 1. Formation of the microstructure of Nd2.sFe13COCl.425B0.o7 s at 800 °C. The starting microstructure consists mainly of free iron and an Nd-rich phase and was obtained by a previous heat treatment at 1150 °C for 15 min.

Ch. Hellwig / Thermomechanical treatment of NdzsFe13CoCl.~25Bo.o75 159

100

90 8 5 0 ° C ~'~ 80

c 7 0

.o 60 u

5 o

4 0 ~ / ' f ± ~ - p h o s e / 2

E 3 O

-S 20

> 10 - p h o s e I 0 . . . . . . . . i . . . . . . , , I . . . . . . . . I

10 100 1000 o n n e o l i n g t i m e [ m l n ]

Fig. 2. Formation of the microstructure of Nd2.sFe13CoCl.42sB0.oTs at 850 °C. The starting microstructure consists mainly of free iron and an Nd-rich phase and was obtained by a previous heat treatment at 1150 °C for 15 min.

100

90 9 0 0 ° C ~' 80 f r e e iron J ~ " I

c ,o : : .9 60 t.) o 50

40 e ~ - p h a s e

~o i_______~__~ I > 10 - 0 p h o s e ., . . . . . . . . , . . . . . . . . ,

10 100 1000 o n n e o l l n g t i m e [ r a i n ]

Fig. 3. Formation of the microstructure of Nd2.sFe13CoC1.425B0.o75 at 900 °C. The starting microstructure consists mainly of free iron and an Nd-rich phase and was obtained by a previous heat treatment at 1150 °C for 15 min.

} ;7

Fig. 5. Micrograph of the Nd2.sFelaCo(C0.gsB0.05)l.5 sample annealed at 850 °C for 1 rain after a previous heat treatment at 1150 °C for 15 min. The sample was polished with diamond particles ( < 1 /sin) and subsequently etched in 10% acetic acid in ethanol. The large grains are supposed to be a-iron whereas the clark areas are due to the Nd-rich phase that dissolves during etching. Only a small amount of intermetallic phases are already formed at the iron/Nd-rich interface.

100 90 9 5 0 ° C "

~' 80 free iron T T /

o,o .2 6o

~, 50

4 0 . "

"6 20 [_ Nd-rich - ~ T ~--------i

> 10 ~ p h o s e , , .., , 0

1 10 100 1000 o n n e o l i n g time [rain]

Fig. 4. Formation of the microstructure of Nd2.sFe13CoC~.4~B0.o75 at 950 °C. The starting microstructure consists mainly of free iron and an Nd-rich phase and was obtained by a previous heat treatment at 1150 °C for 15 min.

samples were cut and characterized in two orientations, parallel and perpendicular to the flow direction of the material. The flow direction is parallel to the defor- mation tool (two parallel plates). An X-ray measurement

Fig. 6. Micrograph of the NdzsFe13Co(C0.95B0.05)1.5 sample annealed at 850 °C for 30 rain after a previous heat treatment at 1150 °C for 15 min. The amount of ¢' phase is grown. A small strip of A phase is visible between the free iron and the 4) phase.

of this parallel orientation will deliver the intensities of crystal planes also parallel to the deformation tool. Micrographs of a parallel orientation show the direction of the applied force perpendicular to the image plane.

Figure 7 shows the diffraction patterns of sample no. 1 after hot deforming in parallel and perpendicular orientation and a calculated pattern. The pattern of the parallel orientation shows preferred orientation. Reflections with high /-indices are stronger than cal- culated (compare especially the 105, 313 and 115 re- flections and the 224 and 410 reflections). Furthermore the intensity of the a-iron reflection is strongly reduced

160 Ch. Hellwig / Thermomechanical treatment of Nd2jFe13CoC l.,25Bo.o75

TABLE 2. Process parameters for the thermomechanical treated samples. The compositions are: Nd2.sFe13CoC1.425B0.075 for No. 1-6 and Nd2.6FelzTCoCi.425B0.075Cuo.3 for No. 7-9. The temperature is 850 °C

No. Thickness (nun)

Before After

Description of the thermomechanical t reatment (according to the heat t reatment at 850 °C)

1 3.5 2.8

2 3.5 1.0

3 3.5 1.0

4 3.2 2.1

5 3.8 2.6

6 3.7 3.1

7 3.8 2.7

8 3.9 2.7

9 3.7 3.2

The whole deformation was made during the first minute

The whole deformation was made during the first minute

The whole deformation was made during the first minute subsequently annealed during 10 min under 75 MPa

Annealed for 2 min then reduction of the thickness of 0.5 mm then each minute 0.2 mm (3rd, 4th, 5th min)

Annealed for 5 min then reduction of the thickness each minute of 0.2 mm (5th, 6th, 7th, 8th, 9th, 10th min)

Annealed for 4 min then reduction of the thickness each minute of 0.1 mm (5th, 6th, 7th, 8th, 9th, 10th min)

Annealed for 5 min then reduction of the thickness each minute of 0.2 mm (5th, 6th, 7th, 8th, 9th, 10th min)

Each minute reduction of the thickness of 0.2 mm (0th, 1st, 2nd, 3rd, 4th, 5th min)

Annealed for 4 min then reduction of the thickness each minute of 0.1 mm (5th, 6th, 7th, 8th, 9th, 10th min)

¢-

35 37

204 214 /

Fig. 7. Observed and calculated diffraction pat terns of sample no. 1 (a) observed intensities of the deformed surface (orientation: parallel), (b) observed intensities of the hole thickness (orientation: perpendicular), (c) calculated intensities of isotropic ¢, phase, x = 1.5418 A.

313

224

a-iron

314 ~ 331.

I I i i I I I I I I I I F I

39 41 43 45 2 0 [° l

; ; 50 ~Jm

in the parallel orientation. The pattern of the per- pendicular orientation shows no preferred orientation. Figure 8 shows the sample no. 1 in perpendicular orientation. Only the microstructure at the surface was affected by the deformation while the inner part of the sample stays unaffected. It is obvious that the X- ray intensities of the parallel orientation were mainly affected by this surface structure while the intensities of the perpendicular orientation were affected mainly by the undeformed structure of the inner sample.

Samples with higher deformation rate (no. 2 and 3) show a completely deformed microstructure (Fig. 9). The great disadvantage of this strong deformation is

Fig. 8. Micrograph of sample no. 1 in perpendicular orientation. Only the surface region was affected by the deformation.

Ch. Hellwig / Thermomechanical treatment of Ndz~Fe13CoCt.425Bo.o75 161

T A B L E 3. The ratios of the intensities/'(105)'/(313) of all measured samples (i.e. hot deformed, annealed in an induction f u r n a c e = i n d l , ind2 and in a tempera ture g rad ien t= t .g r . ) and

the calculated ratio /"(105):1(313) for isotropic Nd2Fe14C. The ori- entat ion is given by II and _1_

Fig. 9. Micrograph of sample no. 2 (orientation: perpendicular) . The whole material was affected by the deformation. The free iron dendrites are flat while a great amount of intergranular Nd- rich phase was pressed out during deformation.

No. Af ter t rea tment After t rea tment and annealing

1(105):](313) II [(105):/(313) 1 1(i05):1(3t3) II 1(105):1(313 ) 2.

1 1.95 0.73 0.91 -- 2 2.15 -- -- -- 3 1.41 -- -- -- 4 1.14 0.53 -- 0.63 5 -- 0.77 -- 0.76 6 -- 0.83 -- 0.70 7 -- 0.62 -- 0.84 8 -- 0.69 -- 0.75 9 - 0.80 - 0.89

t.gr. 2.71 0.32 indl 0.83 0.87 ind2 0.87 0.93 calc 0.67 0.67 0.67 0.67

Fig. 10. Micrograph of sample no. 5. Three different regions are visible (from left to right): inner region with excess iron, an aggregation of Nd-rich phase and the outer region with unaffected microstructure.

the migration of the Nd-rich phase. The Nd-rich phase is not completely solidified at 850 °C and therefore is pressed out of the sample during the deformation. The lack of Nd in the sample leads to undesired excess iron.

The samples No. 4-9 were deformed stepwise in order to affect the whole microstructure without losing the Nd-rich phase. Micrographs (see Fig. 10) show the formation of three radial regions. After an annealing at 850 °C for 3 weeks, the inner region still shows a certain amount of a-iron. A part of the Nd-rich phase was again pressed out and gathered in the second radial region, while the microstructure in the third outer region is still unaffected.

The results of the X-ray measurements are listed in Table 3. The only range without reflections of the ,~

phase was found between d = 2.36 and d = 2.25. There- fore only the ratios of two reflections of the q~ phase, which lie within this range, are given in Table 3. These two reflections are the 313 and the 105 reflection. The 105 reflection will become more intense when the sample shows a preferred orientation of the z axis perpendicular to the measured surface. The intensity of the 313 reflection is expected to stay equal or to become weaker for an oriented sample.

The samples no. indl and ind2 show no significant differences in both orientations and therefore are ex- pected to be isotropic and are used as standard values. In this work ratios 1(105):1(313 ) smaller than 0.75 or greater than 1.0 were assumed to be significant for anisotropy. The calculated relation for Nd2Fe14C is smaller than these standard values due to experimental conditions. According to this standard range only sample no. 4 shows significant anisotropy before and after annealing.

Only pieces of the inner (a-iron rich) region were measured with the moving sample magnetometer. The results are listed in Table 4. The absolute value of the saturation magnetization is strongly affected by the a- iron content and therefore is not discussed. The dif- ferences in Ms of the orientations show a slight preferred orientation in samples no. 4 and 8.

3.3. Using the induction furnace and anneatmg a temperature gradient

The samples annealed in the induction furnace were measured parallel and perpendicular to the applied high frequency magnetic field. Neither the diffraction

162 Ch. Hellwig / Thermomechanical treatment of NdzsFe13CoC l.~zsBo.o75

TABLE 4. The saturation magnetizations Ms of the hot deformed and annealed samples. The orientation is given by II and _1_ with regard to the direction of the material flow during deformation. Only samples with a pronounced hysteresis were measured in both directions

No. After hot deforming and annealing

M~II Ms± (Am 2 kg -1) (AITI 2 kg -I)

1 -- 99 4 130 134 5 -- 149 6 -- 123 7 -- 112 8 121 123 9 85 86

measurements nor the magnetic measurements of these samples show preferred orientation.

The sample annealed in the temperature gradient was cut perpendicular to the applied temperature gra- dient where the temperature was about 850 °C during annealing. Diffraction patterns were taken from this cutting plane (i.e. perpendicular to the temperature gradient) of the low temperature part and from the side plane (i.el parallel to the temperature gradient) of the high temperature part. The results are also listed in Table 3 and show a highly anisotropic microstructure.

4. Discussion

The starting microstructure was obtained by a previous heat treatment at 1150 °C for 15 min and consists mainly of dendrites of free iron with an intergranular Nd-rich phase. The desired tetragonal ~ phase starts forming immediately at a subsequent annealing at around 850 °C at the interfaces between these two phases (see micrograph in Fig. 5). After 15 min, more than 50% of the material consists of the q~ phase. A small amount of the rhombohedral A phase is formed around the iron grains (see micrograph in Fig. 6). The formation of new qb phase can only occur by the diffusion of iron atoms through the already formed intermetallic phases to the Nd-rich phase, therefore long annealing times (several days) are required afterwards to eliminate free iron. At 800 °C the diffusion slows down the formation of ~, while at 900 °C and 950 °C the h phase becomes more stable.

The transformation at 850 °C is fast enough that the application of a thermomechanical treatment seems to be promising. The material can be deformed during the formation of the intermetallic phases. X-ray analysis of the deformed surface of sample no. 1 yields an estimated amount of oriented crystals of 65% of all phase crystals. The responsible mechanism can be ex-

plained as follows. The ~ phase prefers growing in xy direction [22]. Therefore the crystallites formed at the iron/Nd-rich interfaces are flat crystals with the z axis perpendicular to these interfaces. In case of plastic deformation, the crystallites break into platelets and align with the xy plane parallel to the flow direction. The determining factor for preferred orientation in this alloy is therefore the plastic deformation and not the pressure during recrystallization as described for Nd-Fe-B alloys [18]. The plastic deformation yields to the formation of new iron/Nd-rich interfaces and the decomposition of the a-iron is promoted. It is therefore supposed that the reduced intensity of the a-iron re- flection refers to a lower a-iron content and not to a preferred orientation of the free iron. The texture in iron is normally caused by deformation and recrystal- lization and not by deformation alone. The stepwise deformation of the samples to prevent migration of the Nd-rich phase was not successful. Furthermore the samples deformed in this manner showed nearly no preferred orientation. The only sample with significance for preferred orientation due to X-ray analysis and magnetization measurements was deformed for 16% in the beginning in comparison with a maximum 5% for the other samples. It is therefore assumed that a large and continuous deformation is necessary to obtain a preferred orientation. The influence of the amount of already formed ~ phase at the beginning of the de- formation is still unclear. The deformation of sample no. 4 started after 2 min whereas that of sample no. 5 started after 5 min but with different deformation rates. Therefore further experiments are necessary to optimize the period of annealing at 850 °C before deformation. Other possible techniques to prevent the migration of the nearly liquid Nd-rich phase during deformation could be the following:

(1) deformation at temperatures below 850 °C. First experiments are promising whereas the deformation speed has to be chosen very carefully to prevent cracking.

(2) The use of an iron shell (as described in ref. 9 for NdaFe14B compounds). This technique could be combined with previous rheocasting to prevent the formation of an iron dendrite network or

(3) a combination of these techniques. Annealing in a temperature gradient seems also to

be a useful technique to obtain anisotropic bulk material. An explanation of this behaviour cannot be given. A connection with the preferred growing direction of the crystals in the xy plane is supposed. To obtain a tem- perature of 850 °C in a material combined with a temperature gradient, the material at one side must have higher temperatures where more A phase is formed. It is highly expected that this A phase recrystallizes in the • phase at a subsequent annealing at 850 °C with respect to the preferred orientation because the main

Ch. Hellwig / Thermomechanical treatment of NdzsFe13CoC i. 425Boo75 163

process for recrystallization in this microstructure at this tempera ture will be grain growth and not the nucleation of • phase. Fur thermore compositions with a slightly higher boron content could be tested to decrease the A phase formation at higher temperatures.

paring the samples, Mr. K. Mattenberger for magnet- ization measurements and Mr. E. Brtihlmann (Sulzer Medizinaltechnik AG) for assistance with the hot de- formation experiments. This work has been supported by the Swiss National Science Foundation.

5. Conclusions

It is possible to obtain anisotropic bulk material by using the undercooled peritectic reaction y-iron + Nd- rich phase ~ q~. The A ~ q) transformation is not useful for that purpose because of its large temperature range [20]. The present investigations lead to the following statements for a thermomechanical treatment:

(1) the orientation of the crystallites is related to the degree of plastic deformation and not to the applied pressure as reported for NdzFelaB samples [18].

(2) New interfaces between free iron and Nd-rich phase are formed by the breaking of the intermetallic crystallites. Therefore the formation of the 45 phase and the decomposition of the free iron is promoted. This was also observed by Kwon et al. for P r - F e - B - C u samples [23].

(3) The alloy for this kind of thermomechanical t rea tment need not undergo a complete phase transition A--* 4, only the peritectic reaction is necessary.

Therefore the composition can be slightly shifted to the boron-rich side in order for further stabilization of the q~ phase and to allow higher annealing tem- peratures after hot deforming. Nickel can be added for bet ter corrosion behaviour [20].

The problem of the migration of Nd-rich phase out of the material during hot deforming could not be solved. Possible solutions are pointed out. A preferred orientation can also be achieved by heat t reatment in a temperature gradient. Fur ther experiments with mov- ing samples in a zone melting apparatus and/or with samples with a higher boron content are necessary to obtain fully recrystallized anisotropic bulk material.

Acknowledgements

I gratefully thank Professor P. Wachter and Dr. V. Gramlich for their interest, Mrs. M. Krichel for pre-

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