MICROSTRUCTURE EVALUATION OF IRON NITRIDE INTERSTITIAL COMPOUND, AS A CANDIDATE FOR PERMANENT MAGNETIC MATERIAL
by
PARIVASH MORADIFAR
Submitted in partial fulfillment of the requirements
For the degree of Master of Science
Thesis Advisor: Prof. Dr. rer. nat. habil. Frank Ernst
Department of Materials Science and Engineering
CASE WESTERN RESERVE UNIVERSITY
May 2016
! Committee Approval Sheet
CASE WESTERN RESERVE UNIVERSITY
SCHOOL OF GRADUATE STUDIES
We hereby approve the thesis of
PARIVASH MORADIFAR
candidate for the Master of Science*
Committee Chair FRANK ERNST
Committee Member DAVID MATHIESEN
Committee Member PETER D. LAGERLOF
Date of Defense January 2016
!
*We!also!certify!that!written!approval!has!been!obtained!for!any!proprietary!material!contained!therein.!
2! Dedication
This thesis is gratefully dedicated to my beloved parents, Farideh Barati Hezaveh and Mohammad-Esmaeil Moradifar, my beloved sister, Paria Moradifar, for their unconditional love, support, and guidance, without which this thesis would have never been finished.
3! Table of Contents
Table of Contents ...... 4!
List of Figures ...... 6!
List of Tables ...... 9!
Abstract ...... 12!
Chapter 1 -! Introduction ...... 15!
Chapter 2 -! Background ...... 18!
2.1! Permanent Magnet Materials ...... 18!
2.2! Iron-Nitrogen (Fe-N) System ...... 19!
2.2.1! Interstice sites in iron atom (BCC and FCC) structure ...... 20!
2.2.2! Nitrogen austenite (γ- (Fe-N)) ...... 21!
2.2.3! Nitrogen-martensite (α') ...... 22!
2.2.4! γ'-Fe4N ...... 25!
2.2.5! Ordered martensitic nitride (α''-Fe16N2) ...... 26!
2.3! α''-Fe16N2, a phase with a possible giant magnetization ...... 28!
2.4! TEM Sample Preparation of Powder Particles ...... 34!
2.5! Purpose of Study ...... 36!
Chapter 3 -! Materials and Experimental Methods ...... 38!
3.1! Materials and Nitriding Process ...... 38!
3.2! Characterization Technique ...... 41!
3.2.1! X-Ray Diffraction (XRD) ...... 41!
3.2.2! Vibrating Sample Magnetometer (VSM) ...... 43!
3.2.3! Scanning Electron Microscope (SEM) ...... 43!
3.2.4! Transmission Electron Microscope (TEM) ...... 45!
4! Chapter 4 -! Results ...... 46!
4.1! XRD Results ...... 46!
4.2! VSM Results ...... 50!
4.3! TEM Sample Preparation Results ...... 52!
4.4! TEM Results ...... 56!
4.4.1! TEM observation of quenched sample ...... 56!
4.4.2! TEM observations of cryomilled sample ...... 67!
4.4.3! TEM observations heat-treated sample ...... 75!
Chapter 5 -! Discussion ...... 90!
5.1! XRD Results ...... 90!
5.2! VSM Results ...... 92!
5.3! TEM Sample Preparation Results ...... 94!
5.4! TEM Results ...... 95!
5.4.1! Quenched sample ...... 95!
5.4.2! Cryomilled sample ...... 97!
5.4.3! Heat-treated sample ...... 98!
Chapter 6 -! Conclusion ...... 102!
Chapter 7 -! Future Work ...... 106!
5! List of Figures
Figure 1: Applications of permanent magnets ...... 15!
Figure 2: History of permanent magnetic materials [1] ...... 16!
Figure 3: China rare earth exports pricing ...... 17!
Figure 4: Magnetization versus magnetic field strength (M vs. H) curve [8] ...... 19!
Figure 5: Iron-Nitrogen system [9] ...... 20!
Figure 6: Interstitial sites in (a) FCC and (b) BCC iron [10] ...... 21!
Figure 7: Structure of nitrogen-austenite ...... 22!
Figure 8: Nitrogen-martensite (α') with BCT structure (a=b≠c, α=β=γ=90°) [10] ...... 24!
Figure 9: Lattice parameters variation of iron-nitrogen martensite (α') and iron-carbon martensite with interstitial content [16] ...... 25!
Figure 10: γ'-Fe4N structure [17] ...... 26!
Figure 11: Structure of α''-Fe16N2 [21] ...... 27!
Figure 12: (a) Nitriding reactor, (b) Schematic of nitriding reactor by D.Matthiesen ...... 39!
Figure 13: Schematic of producing bulk samples of α''-Fe16N2 ...... 40!
Figure 14: Cryomilling machine ...... 41!
Figure 15: Lift-out protocol of a slice of a single particle ...... 44!
Figure 16: XRD diffractogram of AHC 100.29 as-received and simulated diffractogram of α- iron ...... 46!
Figure 17: simulated XRD diffractograms of γ-iron, α'-Fe8N, and α''-Fe16N2 ...... 47!
Figure 18: XRD patterns of quenched sample, cryomilled sample and heat-treated sample ...... 49!
Figure 19: Superimposed XRD diffractograms of cryomilled sample and heat-treated sample .. 50!
Figure 20: VSM curves of four samples of (a) As received sample (b) Quenched sample (c) Cryomilled sample (d) Heat-treated sample ...... 51!
6! Figure 21: SEM image of cross-sectional embedded particle powders in the epoxy resin (Durcupan) ...... 53!
Figure 22: SEM image of cross-sectional embedded particle powders in commercial solder ..... 54!
Figure 23: XEDS map of cross-sectional embedded particle powders in commercial solder ...... 54!
Figure 24: STEM image of cross-sectional embedded particle powders in commercial solder (a) low magnification (b) higher magnification of marked area ...... 55!
Figure 25: XEDS map of cross-sectional embedded particle powders in commercial solder ...... 56!
Figure 26: (a) low mag STEM image of nitride sample at CL: 120 mm (b) magnified STEM image from the bottom of main needle at CL: 250 mm (c) magnified STEM image from the tip of main (central) needle at CL: 250 mm ...... 58!
Figure 27: Bright field images from bottom, middle and tip of the area containing the big needle in low mag STEM image ...... 59!
Figure 28: (a) electron diffraction pattern corresponding to bottom part of the needle (area marked 1) (b) shows the grey level adjusted of central area (c) shows the simulated electron diffraction pattern for [111] zone axis of α''-Fe16N2 (d) solved electron diffraction pattern for α''-Fe16N2 phase at [111] zone axis by using JEMS software (e) electron diffraction pattern corresponding to top part of the needle (f) grey level adjusted of central area of part (e) .... 62!
Figure 29: (a) shows an obtained diffraction pattern from area 2 marked on Figure 27(b) shows the simulated diffraction pattern for the [110] zone axis of γ-Fe phase (c) solved electron diffraction pattern for the [110] zone axis of γ-Fe phase by using JEMS software ...... 63!
Figure 30: Dark field image obtained from one of the super-lattice reflections belonging to {112} planes ...... 64!
Figure 31: (a) Bright field image of top right part of sample including smaller needles and matrix (b) electron diffraction pattern of matrix (c) electron diffraction pattern of marked area containing small needle and surrounded matrix (d) simulated electron diffraction pattern of α'-Fe8N for [100] zone axis ...... 66!
Figure 32: (a) First and (b) second STEM images of cryomilled sample at CL:250 mm (c) simulated electron diffraction pattern of γ-iron for [110] zone axis (d) electron diffraction pattern of marked area in b (e) bright field image of corresponding area in part (a) ...... 69!
Figure 33: (a) martensitic lath in cryomilled sample (b) electron diffraction pattern from marked area (c) simulated diffraction pattern of α'-Fe8N for the [113] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [113] zone axis by using JEMS software ...... 71!
7! Figure 34: (a) shows a TEM image containing the area the electron diffraction pattern obtained (b) obtained electron diffraction pattern (c) simulated diffraction pattern of α'-Fe8N for [111] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [111] zone axis by using JEMS software ...... 73!
Figure 35: (a) small laths and the area that the electron diffraction pattern obtained (b) electron diffraction pattern of marked area (c) simulated diffraction pattern of α-iron for [100] zone axis (d) solved electron diffraction pattern of α-iron for [100] zone axis by using JEMS software ...... 75!
Figure 36: STEM images of heat-treated sample (a) and (b) obtained at CL: 120 mm (c) obtained at CL: 80 mm ...... 76!
Figure 37: (a) (b) (c) (d) bright field (BF) images of three sets of needles in tempered sample in different magnifications ...... 78!
Figure 38: (a) TEM image after tilting (b) electron diffraction pattern of marked area (c) simulated electron diffraction pattern of α'-Fe8N for [111] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [111] zone axis by using JEMS software ...... 80!
Figure 39: (a) TEM image of tilted sample, marked area shows the corresponding area of obtained diffraction pattern (b) electron diffraction pattern of marked area (c) grey level adjusted electron diffraction pattern to reveal superlattice reflections (d) simulated electron diffraction pattern of α''-Fe16N2 for [100] zone axis (e) solved electron diffraction pattern of α''-Fe16N2 for [100] zone axis zone axis by using JEMS software ...... 83!
Figure 40: Dark field images corresponding to (a) and (b) {011} planes (c) {002} planes in α''- Fe16N2 electron diffraction pattern at [100] zone axis ...... 84!
Figure 41: (a) area of obtained electron diffraction pattern (b) corresponding electron diffraction pattern of marked area (c) simulated electron diffraction pattern of α''-Fe16N2 for the [103] zone axis (d) solved electron diffraction pattern of α''-Fe16N2 for [103] zone axis by using JEMS software ...... 86!
Figure 42: dark field image corresponding to spot from the {301} planes in α''-Fe16N2 electron diffraction pattern at the [103] zone axis ...... 87!
Figure 43: (a) (b) virtual TEM images and DPs of area near to the surface of needle (c) (d) virtual TEM images and DPS of central area of needles (e) TEM image and DP of bright area of matrix ...... 88!
Figure 44: (a) (b) (c) virtual dark field images corresponding to superlattice reflections marked on diffraction pattern (d) virtual dark field image corresponding to marked main spot on diffraction pattern ...... 89!
8! List of Tables
Table I: Processing conditions of samples ...... 41!
Table II: Magnetic properties measurement of the samples ...... 52!
9! Acknowledgements
I would like to express my sincere gratitude and appreciation to my supervisor, Professor
Frank Ernst, for his valuable advice, guidance, mentorship, and full support throughout my master’s study at Case Western Reserve University. I would also like to express my appreciation to my thesis committee members, Prof. David Matthiesen and Prof. Peter D. Lagerlof for their participation, invaluable advice, thoughtful comments and suggestions to increase the quality of this work.
I should also acknowledge contributions of Prof. David Matthiesen, Prof. Matthew Willard,
Zhiyao Feng and Song Lan in conducting nitridation experiments, VSM measurements and their thoughtful advice and help. This work was supported by Advanced Research Projects Agency-
Energy (ARPA-E).
I would also thank SCSAM engineers, Dr. Amir Avishai, Dr. Wayne Jennings, Dr. Jonathan
Cowen, Nan Avishai and specifically Dr. Danqi Wang for his help and advice in TEM part of this project. I would also like to thank my previous and current colleagues in Prof. Frank Ernst research group, Dr. Zahra Vashaei, Dr. Thanga, Xun, Anna, Qinong, Zhen and Amirali for their help and useful discussions throughout these years.
My deepest gratitude and indebtedness goes to my dear parents, Farideh and Mohammad-
Esmaeil and my sister, Paria, for all the sacrifices they made to make my life better and more successful and a beautiful place to live. Without their support, encouragement and understanding it would have been impossible for me to finish this work.
10!
11! Microstructure Evaluation of Iron Nitride Interstitial Compound, as a Candidate for Permanent Magnetic Material
Abstract
by
PARIVASH MORADIFAR
Alternative technologies are developing to replace rare-earth permanent magnets because of high costs and limited supply. Rare-earth permanent magnets currently are used in electric-vehicle motors and wind turbines because of their high saturation magnetization and high coercivity.
Iron-nitrogen magnets (α''-Fe16N2) with possible giant magnetization, can be a promising candidate for replacing rare-earth permanent magnets.
In this study, during nitridation experiments which have been carried out by Z.Feng according to
Jack’s route, nitrogen austenite formed by nitriding the AHC 100.29 (α-Fe) powders with a gas mixture of 0.11NH3/ 0.89 H2 for 3600s (1h) at 923.15K (650ºC) in a designed nitriding reactor.
The nitrided sample containing nitrogen austenite is cryomilled by ball milling the powder at liquid nitrogen temperature for 600s (10 min) at 30 Hz frequency in order for the transformation from nitrogen-austenite to nitrogen martensite to take place, and finally the sample is heat-treated for 7200s (2h) at 403.15K (130ºC) in order to form an ordered nitrogen martensite (α''-Fe16N2) phase with possible giant magnetization. Various characterization techniques, such as transmission electron microscopy (TEM), scanning electron microscopy (SEM), X-ray diffraction (XRD) analysis and vibrating sample magnetometer (VSM) analysis by S.Lan, have been applied to precisely characterize these samples.
12! Nitrogen austenite with the lattice expansion of 2.3% and the nitrogen content of (9.8 ± 0.4) at.% formed by nitriding the AHC 100.29 (α-Fe) powders. Formation of a small fraction of coherent needle-shaped nitrogen martensite precipitations with possible preferred orientation
relationship of [100]α'" [101]γ and (011)α'" (111)γ to the expanded austenite matrix, matched with a
Nishiyama-Wasserman (N-W) orientation relationship and local rearrangement of nitrogen atoms in small parts of formed precipitations have been observed. Around 90% of the nitrogen austenite transformed into nitrogen martensite with lath-shaped morphology after cryomilling, while around 9% of the nitrogen austenite remained as retained austenite after cryomilling with the same lattice parameter of its parent expanded austenite which also had lath-shaped morphology, possibly formed because complex shear interaction during martensitic transformation with a high density of dislocations around them has been observed.
Formation of α''-Fe16N2 precipitations in a heat-treated sample was also confirmed with
VSM. Mass magnetization at 1.5 × 107/4π (A/m) (15 kOe) increased from 158.4 A.m2/kg for the cryomilled sample to 202.5 A.m2/kg for the heat-treated sample can confirm the formation of ordered nitrogen martensite (α''-Fe16N2) after heat-treating the sample from disordered nitrogen martensite (α') from the magnetic point of view.
Formation of coherent α''-Fe16N2 in α' matrix in a heat-treated sample was confirmed with
TEM observation. Electron diffraction patterns obtained from needle-shaped precipitations after heat-treating the sample in different zone axes of [100] and [103] clearly have shown the formation of α''-Fe16N2 in the heat-treated sample. Three sets of surface martensitic needles with different sizes composed of a central untransformed region of nitrogen martensite and a transformed region of nitrogen martensite to α''-Fe16N2 confirming the partial transformation of 13! α' to α''-Fe16N2 precipitations, have been observed after heat-treating the sample. This partial transformation of disordered nitrogen martensite (α'-Fe8N) into ordered nitrogen martensite (α''-
Fe16N2) can explain the discrepancy in magnetic results for α''-Fe16N2 among results obtained by different research groups in the last forty or more years.
14! Chapter 1 -! Introduction
Permanent magnets play an integral role in modern technology and have a wide range of applications from electromechanical machines to wind turbines, energy storage applications, electrical vehicles, electronic devices and even appliances, as shown in Figure 1.
http://www.cndailymag.com
Figure 1: Applications of permanent magnets
Figure 2 shows the history of permanent magnet materials. However, the history of permanent magnet materials goes back to the 1930’s when Alnico, an iron based alloy containing aluminum (Al), nickel (Ni) and cobalt (Co) which has a high Curie temperature of around
1123.15 K (850ºC) was introduced. Later, in the 1950’s, hard ferrite magnets (BaFe12O19 or
SrFe12O19) were introduced by researchers at Phillips Laboratory. There was a revolution in permanent magnets in the 1970’s by the introduction of rare earth-cobalt (Sm-Co base) alloys.
For many years, SmCo5 was very common due to its large energy product. In the 1980’s,
15! neodymium-iron-boron (Nd-Fe-B) based permanent magnets, which had a high energy product were invented. However, it has relatively low Curie temperature of 585.15K (312ºC) [1-4].
Figure 2: History of permanent magnetic materials [1]
For current high performance motors and generators, Nd-Fe-B permanent magnets are widely used with added dysprosium (Dy) and cobalt (Co) to modify its temperature characteristics for high temperatures applications. However, because the mineral resources for rare earth elements such as Nd, Dy, Sm are found in distinct regions of the world, there is a major problem with mineral resources of rare earth elements especially for Dy on an international scale [5]. China controls around 97% of the rare earth elements related markets worldwide and is the largest supplier of rare earth elements. China has put quotas on exports of rare earth elements and increased its price from 2010, as shown in Figure 3 [6].
16! Source: China Customs
Figure 3: China rare earth exports pricing
Iron nitrogen permanent magnets can be a promising candidate for future rare earth free permanent magnetic materials. The iron nitrogen permanent magnets are made out of iron and nitrogen, two abundant and cheap elements rather than requiring any rare earth elements. Also, the theoretical limit of magnetic properties such as energy product (BH)max in iron nitrogen permanent magnets can reach almost twice that of Nd2Fe14B magnets [5]. David Matthiesen, associate professor of materials science and engineering at Case Western Reserve University, said “A costing model puts neodymium–iron–boron used now at $60 per kilogram and our material (iron nitrogen permanent magnets) at about $10 per kilogram.”
17! Chapter 2 -! Background
2.1! Permanent Magnet Materials
Permanent magnets are a group of magnetic materials which can be magnetized by an external magnetic field. A permanent magnet should be able to not only maintain its magnetic behavior after removal of external magnetic field but should also remain magnetic in the presence of an opposing magnetic field. Therefore, theoretically a high performance permanent magnet should have a high coercive force Hc to stand with demagnetization. Coercivity measures the ability of a permanent magnet to stand in the presence of an opposing magnetic field. Since permanent magnets are subjected to relatively high temperatures in some applications, such as modern motors for HEV and EV applications, they should be able to have stable magnetic properties in relatively high working temperatures too [7]. Therefore, a desirable permanent magnet should not only have a high coercivity but also should have a high Curie temperature Tc. The Curie temperature, Tc, is a temperature at which a sharp change occurs in magnetic behavior of a material. For a permanent magnet material, permanent magnetism will change to induced magnetism at this temperature. That is why it is of utmost importance that a permanent magnet be able to keep its magnetic order at relatively high temperatures. In addition, a permanent magnet can be considered as a source of a magnetic field. Therefore, it requires a high remanence, Mr, in order to produce a large magnetic field. Remanence Mr, measures the remained magnetization after the external magnetic field is removed [7]. Figure 4 shows a magnetization versus magnetic field strength (M vs. H) curve.
18!
Figure 4: Magnetization versus magnetic field strength (M vs. H) curve [8]
2.2! Iron-Nitrogen (Fe-N) System
Figure 5 shows the Fe-N phase diagram at 0.1 MPa of hydrostatic pressure. Different phases of iron and equilibrium iron-nitride phases can be seen in this phase diagram. Different phases of iron include: (1) ferrite, a BCC solid solution with a space group of Im3m, α-Fe or δ-Fe, with a stability range of below 1185.15 K (912ºC) and above 1667.15 K (1394ºC) respectively, with maximum nitrogen solubility of 0.3 at. %; (2) austenite (γ-Fe), a FCC solid solution with a space group of Fm3m, with maximum nitrogen solubility of 10.3 at.%. Different iron-nitride equilibrium phases, which can be seen in this diagram include: (1) Fe4N or γ', a nitride phase with FCC structure which appears at around 20 at. % N; (2) Fe2-3N or ε, a nitride phase with close-packed hexagonal structure, which appears from about 15 at.% of N to at least 33 at.% of
N; (3) Fe2N, a nitride phase with orthorhombic structure, which appears from about 33 at.% N to an undefined upper composition [9].
19!
Figure 5: Iron-Nitrogen system [9]
There are also two non-equilibrium nitride phases at 0.1 MPa of hydrostatic pressure, (1)
Fe8N, α', a disordered martensite and (2) Fe16N2, α'', an ordered martensite, both with body centered tetragonal structures [9].
2.2.1! Interstice sites in iron atom (BCC and FCC) structure
There are two types of interstice sites in iron; tetrahedron sites and octahedron sites. A tetrahedron site has four metal atom neighbors, while an octahedron site has six metal atom neighbors as shown in Figure 6. Nitrogen, as an interstitial atom, always occupies the octahedron sites in both BCC and FCC iron. As shown in Figure 6(a), an octahedron site in BCC is distorted
20! in the one direction. Therefore, placing an interstitial atom such as nitrogen in this case will cause a large anisotropic strain in neighboring matrix. That is why the solubility of nitrogen in α- iron is very limited [10]. While an octahedron site in FCC iron, as shown in Figure 6(b), is a symmetric site. Therefore, an isotropic expansion occurs by insertion of an interstitial atom, such as nitrogen, in an octahedral interstice of FCC iron [10].
Figure 6: Interstitial sites in (a) FCC and (b) BCC iron [10]
2.2.2! Nitrogen austenite (γ- (Fe-N))
In nitrogen austenite, iron atoms have an FCC arrangement and up to one in ten of octahedron sites are occupied by nitrogen randomly without any specific ordering. Nitrogen austenite exists at above 863.15 K (590ºC). Figure 7 shows the structure of nitrogen austenite.
Nitrogen sites are at the center of the FCC structure and mid-points of the edges, while iron 21! atoms have an FCC arrangement [10-12]. The lattice parameters of nitrogen austenite also have a relationship with nitrogen concentration. According to the literature, the relationship between lattice parameters and nitrogen concentration in nitrogen austenite is given according to the equation below for 0 to 11 at.% N [9, 10, 13]:
γ γ γ (1) a = a0 +α ⋅ X N
γ γ a is the lattice parameter of nitrogen austenite after nitridation, a0 is the lattice parameter of elemental austenite γ-Fe without nitrogen and it’s equal to 0.3572 nm in above equation. α is an
-4 γ empirical value and is equal to 7.8×10 nm/ at.% according to Wreidt equation. X N is the nitrogen mole fraction in at.%.
Figure 7: Structure of nitrogen-austenite
2.2.3! Nitrogen-martensite (α')
22! By rapid quenching of nitrogen-austenite from above 873.15K (600ºC), nitrogen martensite
(α') forms through a martensitic transformation [10, 14]. It is accepted that nitrogen martensite has the same composition of its parent austenite with a maximum of one in ten occupied interstices with nitrogen atoms at two positions of c-edge midpoints and at the centers of C faces which may also be retained in the final nitrogen martensite as well. At a low enough reaction temperature, interstitial atoms do not have enough time during the martensitic transformation to make a movement from the occupied sites in austenite to the additional octahedral sites created at the midpoints of a-edges and b-edges, and at centers of A faces and B faces. Therefore, these additional sites are never being occupied. As a result, a nitrogen supersaturated ferrite will form in which the expanded axes of all occupied interstices are pointing in the same direction [10].
That is why nitrogen martensite (α') has a body centered tetragonal (BCT) structure (a=b≠c,
α=β=γ=90°) instead of body centered cubic (BCC) structure [10]. In Figure 8, a structure of nitrogen martensite (α') and possible sites for both nitrogen and iron are shown. The octahedral
& 1 # sites in this structure are located at the midpoints of c-edges corresponding to $0,0, ! points % 2 "
& 1 1 # and also at the centers of C faces corresponding to $ , ,0! points [11]. It is also worth % 2 2 " mentioning that Jack’s 1973 paper stated this structure is an average structure of nitrogen martensite because there are some units distorted by the presence of nitrogen, while there are other units, there are no occupied interstices in them and their dimensions are almost same as α-
Fe with BCC structure [10].
23!
Figure 8:Nitrogen-martensite (α') with BCT structure (a=b≠c, α=β=γ=90°) [10]
The term of tetragonality in an iron martensite phase stands for mean anisotropic distortion in its lattice. Tetragonality (c/a) increases with increasing interstitial content according to the equation below [9]:
a = a0 – α (XN/XFe) (2)
c/a = c0/a0 + β (XN/XFe) (3)
Where a and c are the lattice parameters of nitrogen martensite (α'). c0 is equal to a0 and c0/a0 is equal to 1 primarily before any distortion occurs in nitrogen martensite. a0 is in nm and it is equal to 0.28664. α is an empirical value and is equal to 0.018 nm. β is a dimensionless parameter and is equal to 0.91 according to Wriedt equation. XN and XFe denote concentrations of
N and Fe and their unit is in at.%. Figure 9 shows the lattice parameter variations by increasing the concentration of interstitial atoms [15].
24!
Figure 9: Lattice parameters variation of iron-nitrogen martensite (α') and iron-carbon martensite with interstitial content [16]
2.2.4! γ'-Fe4N
The structure of γ'-Fe4N is shown in Figure 10. As is shown, in this nitride phase, iron atoms have an FCC structure which is the same as nitrogen-austenite, while nitrogen interstitial atoms fill one in four of the octahedral interstices in a fully ordered manner. The γ'-Fe4N has a limited range of homogeneity and it is stable only below 953.15K (680ºC) [10].
25!
Figure 10: γ'-Fe4N structure [17]
2.2.5! Ordered martensitic nitride (α''-Fe16N2)
In 1951Jack first stated that [18] tempering of α'-Fe8N phase at temperatures of about
370 K-420 K will result in the ordering of nitrogen atoms. As a result of nitrogen ordering in
α' phase, α''-Fe16N2 as an ordered martensitic phase will be formed [19]. The crystal structure
of α''-Fe16N2 is shown in Figure 11. α''- Fe16N2 has a body centered tetragonal structure with
an I4/mmm space group. α''- Fe16N2 consists of 8, [2× 2 × 2], α-iron distorted unit cells.
Nitrogen atoms occupy 2 of 48 octahedral interstices in a completely ordered manner. The
lattice-parameters of its unit-cell are a = 0.572 nm, c = 0.629 nm with axial ratio of c/a =
1.10. 16 Fe atoms in this structure occupy three different positions of 8h (x = 0.250) at
& 1 1 1 # & 1 1 1 # ±(x, x,0), ± $ + x, x, !, ±(x,− x,0), ± $ + x, − x, !, 4e (z = 0.3125) at (0,0,± z), % 2 2 2 " % 2 2 2 "
& 1 1 1 # & 1 1 # & 1 1 # & 1 3 # & 1 3 # $ , , ± z! and 4d at $0, , !, $ , 0 , !, $0, , !, $ , 0 , !, while the % 2 2 2 " % 2 4 " % 2 4 " % 2 4 " % 2 4 "
26! & 1 1 1 # nitrogen atoms occupy the 2a positions at (0, 0 , 0) and $ , , ! [11, 20][16]. The % 2 2 2 " limited solubility of nitrogen with a maximum of 10.3 at.% in γ-Fe, means that it is impossible to obtain a pure ordered nitrogen martensite phase of (α''-Fe16N2) (11.1 at.% of N) by Jack’s method [19]. By tempering, not only does a chemical reaction occurs (α' to α and
α'') but also a disorder-order transformation occurs. Jack also stated that no α''- Fe16N2 phase forms initially by tempering α'-nitrogen martensite. Instead martensitic regions with higher concentrations of nitrogen form gradually. By increasing the nitrogen concentration, the tetragonality will also increase until it reaches the tetragonality of α''-Fe16N2, corresponding to α' composition of Fe8N. Then, the disorder-order transformation happens which will result in forming α''-Fe16N2 from α'-Fe8N [14].
Figure 11: Structure of α''-Fe16N2 [21] 27! 2.3! α''-Fe16N2, a phase with a possible giant magnetization
α''-Fe16N2 as a metastable, ordered martensitic phase with a BCT structure was discovered for the first time by Jack in 1951 via the X-ray diffraction (XRD) method [11, 22]. The magnetization property of α''- Fe16N2 remained unstudied until 1972. In 1972, Kim and Minuro
Takahashi at Tohoku University reported an ultra-high saturation magnetization in Fe-N polycrystalline films which have been prepared by reactive vapor deposition in atmosphere of low N2 gas pressures for the first time. The polarization of these films reached 2.64 T. These films consist of a mixture of α-Fe and α''-Fe16N2 phases. They have reported that the high saturation magnetization of these Fe-N films is due to the formation of the α''-Fe16N2 phase. The inferred saturation magnetization for α''-Fe16N2 was 2.9 µB per Fe atom [14, 23, 24]. After that, different research groups worked on the α''-Fe16N2 phase as a phase with possible giant magnetization. Different research groups have used a variety of methods, such as ion- implantation, molecular beam epitaxy (MBE), radio frequency (RF) sputtering, face target sputtering (FTS), reactive plasma sputtering, and reactive plasma evaporation to produce thin films with the highest volume fraction of α''-Fe16N2 that will be reviewed in more detail below.
Three other research groups also tried to produce α''-Fe16N2 in bulk form by using Jack’s route
[18] for producing α''-Fe16N2 in bulk form.
Nakajima and his co-workers prepared thin films by using the ion-implantation method.
High-dose nitrogen ions implanted into the single crystalline iron films until it reaches the proper stoichiometric concentration for α''-Fe16N2. They have reported that the final film consists of α-
Fe, mostly α'-nitrogen martensite and a very small amount of α''-Fe16N2 according to XRD data.
28! Also, the magnetization of primary α-Fe film increased by 13.4% after high-dose nitrogen implantation which is related to the formation of nitrogen martensite (α') and very small volume fraction of α''-Fe16N2. Finally, they have reported that the magnetization of nitrogen martensite, including both α'- nitrogen martensite and α''-Fe16N2, is about 2.57 T [25].
Sugita and his group at Hitachi Laboratories tried to prepare a single-crystal of α''-Fe16N2 film in (001) direction on a (001) In0.2Ga0.8As and (001) GaAs substrate via the molecular beam epitaxy (MBE) method by evaporating Fe in a mixed gas atmosphere of N2 and NH3. The measured total magnetization for these films was 2.9 T on average. The proposed magnetization was 3.2µB per Fe atom. They have also reported that the saturation magnetization of α''-Fe16N2 is temperature dependent [26-30].
Doyle and his co-workers produced single-layer Fe-N films on a glass substrate by using the radio frequency deposition (RF) method to deposit a single layer in an atmosphere of N2 and Ar for different flow rates of N2. For an optimum flow rate of N2, α''-Fe16N2 has been formed in a mixture with α-Fe and the total magnetization reached its maximum of 2.4 T, which corresponds to 3.2 µB per Fe atom. They also confirmed the existence of α''-Fe16N2 not only by XRD but also with TEM [31].
Sun et al. also tried to produce single crystal α''-Fe16N2 thin films by using face target sputtering on a NaCl substrate in an atmosphere of N2 and Ar. They also stated that the substrate temperature is a key factor in forming α''-Fe16N2. They have used both XRD and TEM for phase identification and they also confirmed the initially proposed structure by Jack in 1951 [18]. The
29! measured total magnetization for these films was 2.87 T, corresponding to 3.2 µB per Fe atom which they believe is in agreement with the results of Sugita et al. and Kim et al [14, 23, 26, 32].
In 1996, Brewer et al. used the reactive sputtering method followed by annealing to form α''-
Fe16N2 in produced thin films. They formed α'-Fe8N on Si (001) substrate. According to their analysis, by using XRD technique, after annealing, these films consist of 54% of α' and 46% of
α'' phase. The total magnetization for these films was 2.23 T. The proposed value was 2.4 µB per
Fe atom [14, 33].
Migoku Takahashi and his group produced thin films by reactive plasma sputtering with a thickness of 0.03µm to 1 µm and by reactive plasma evaporation with a thickness of 0.2 µm on
MgO substrate. Thin films prepared by plasma sputtering showed magnetization of 218 A.m2/kg and ones prepared by plasma evaporation showed magnetization of 235 A.m2/kg. These values are much smaller than the initially reported giant magnetization of 2.9T (310 A.m2/kg) for α''-
Fe16N2 [23, 34, 35].
Wallace et al., [36, 37] Metzger et al., [38, 39]and Coey et al., [17] in three different groups, tried to produce pure α''-Fe16N2 in bulk form by using Jack’s method of γ→ α'→ α'' in 1951
[18]. Wallace et al. and Metzger et al., both produced α''-Fe16N2 phase as a mixture with α' and γ
phase and the proposed magnetization for α''-Fe16N2 was 2.4 µB per Fe atom and 2.9 µB per Fe atom respectively, while Coey and his co-workers produced α''-Fe16N2 as a mixture with α', γ' phase. The inferred value by Coey’s group was 2.3-2.6 µB per Fe atom at T~ 0 K.
30! Unless different research groups worked on α''-Fe16N2 as a material with possible giant magnetization, researchers obtained variable results of 240 A.m2/kg to 315 A.m2/kg for the magnetic moment of α''-Fe16N2 and inconclusive [24]. Therefore, due to these variable results on the saturation magnetization of α''-Fe16N2, the magnetism has remained a mystery for more than
40 years [42].
There are also discrepancies among experiments and theoretical calculations [40]. The existence of a giant magnetization in α''-Fe16N2 with the structure proposed by Jack in 1951[18] is not confirmed by band theory or the electronic structure calculations reported by some research groups [21, 41] and there is a discrepancy between their results and other results, such as inferred value for giant magnetization of α''-Fe16N2 by Hitachi’s group [26] [14]. In 1994,
Coey did electronic structure calculation and reported that the magnetization of α''-Fe16N2 is close to α-Fe (2.2 µB per Fe atom) and there is no giant magnetization [21] and Lai et al. also reported that the magnetization of α''-Fe16N2 can be slightly higher (about 10%) than α-Fe [42].
In 2000, Takahashi pointed out that this variable proposed results by variable research groups for magnetization of α''-Fe16N2 can have three different sources. First, there can be errors related to phase identification of α''-Fe16N2 [24]. Most research groups used XRD for phase identification and even quantification of volume fraction, while there is an overlap between the most diffraction peaks of α'' and peaks of α and γ phase [20]. Furthermore, in XRD analysis, only a surface region of powder is probed. Therefore, it is not a precise method for the purpose of volume fraction measurements and phase identification. The second source of errors in measurements, Takahashi stated concerned thin films. There is no reliable method for fixing α''-
31! Fe16N2 phase in a whole film. Third, the lack of accuracy of magnetization of other iron nitride phases measured before are used for determining of α''-Fe16N2 in cases which α''-Fe16N2 exist in a mixture with other iron and iron-nitride phases [24].
There is very little structural information reported on iron-nitrides and α''- Fe16N2 in terms of transmission electron microscope (TEM) work. Previous works primarily concentrate on magnetic properties. less works has been done on structural analysis. Even among these groups, most of them used X-ray diffraction (XRD) analysis for structural analysis [40].
In 1997 Sun et al., did some structural analysis of XRD and TEM on single crystal films of
α''-Fe16N2 produced by facing target sputtering (FTS) on a NaCl (001) single crystals by using two pairs of iron targets and using Ar as sputtering gas and N2 as reactive gas. In addition to
XRD, they also obtained electron-diffraction patterns of this phase in which some of them confirmed the primary proposed structure with Jack (Jack-1) and some of them did not confirm this structure [40].
Hayes and his coworkers also did in situ experiments to study the kinetics and dissolution of
α''-Fe16N2 in thin foils of Fe-based alloy containing oxygen and phosphorous with 0.05 wt% of nitrogen [43]. Xiong et al. also studied the kinetics of α''-Fe16N2 precipitation obtained by aging of supersaturated nitrogen-ferrite. They reported that there is a two stage process for precipitation of α''-Fe16N2 including coarsening and coarsening stabilization. They also reported disc shape morphology for α''-Fe16N2 precipitations which have been formed as islands in a ferrite matrix.
In the second stage, the number of islands decreased and smaller islands dissolved, which is consistent with the Gibbs-Thomson effect [44].
32! Watanabe and his coworkers studied the microstructure of α''-Fe16N2 in terms of the effect of applied stress on formation of α''-Fe16N2 in Fe-N alloys by scanning electron microscope (SEM).
In this study, they also used a magnetic torque meter to measure the magnetic anisotropy. They reported that the applied stress can cause a preferential nucleation of α''-Fe16N2 by their observation via SEM. They also approved this result by measuring magnetic anisotropy. They categorized α''-Fe16N2 precipitation into groups of premature and mature precipitations from a magnetic point of view, but they did not do any further structural analyses to find out the differences between these two groups of precipitations from a structural point of view [45].
Tanaka et al. also did a structural analysis via TEM on iron-martensitic steels. They reported a decomposition of α'-nitrogen martensite into α''-Fe16N2 and low nitrogen martensite α'1 occurring during tempering of α'-nitrogen martensite. They also stated that α''-Fe16N2
precipitations have a plate-like appearance while α1' have a mosaic-like appearance which have parallel directions. They also noticed a new kind of nitride phase, which they call α'''-Fe16N with a base-center orthorhombic crystal structure. According to their report this phase rarely forms coherently into α''-Fe16N2. They also did a structural refinement for α''-Fe16N2, stating that the parameters for 4e and 8h positions in this structure do not match Jack’s earlier model [20].
In 2000, Li and his group also observed α''-Fe16N2 and γ'-Fe4N precipitations in the diffusion layer of iron-nitrided samples which had been nitride in an ion-nitriding furnace. They reported
γ'-Fe4N precipitations as long rods and α''-Fe16N2 precipitations as thin plates in their observations. They also observed electron diffraction patterns of α''-Fe16N2 in different directions. They also studied the orientation relationship between α''-Fe16N2 precipitations and
33! α-Fe the surrounding matrix [18]. In 2001, they studied microstructure of α''-Fe16N2 specifically and also the transformation of this phase to γ'-Fe4N phase. They reported a ribbon-like appearance this time for α''-Fe16N2 phase despite a disc-shape appearance or an undulated lath- shaped feature [46] which has been reported by other research groups, and thin plate-like features that they reported previously [17]. Other researchers, such as Hinojosa et al. tried to find the relationship between α''-Fe16N2 precipitation formation and hardness [47].
2.4! TEM Sample Preparation of Powder Particles
Different methods have been proposed and applied for preparing TEM specimens out of powder particles by different research groups. One method is reducing the powder particles’ size until the appropriate particle thickness is reached for TEM observation, followed with making a suspension from reduced size powders in a volatile liquid and leaving a small drop of this suspension on a carbon grid to become dry and ready for observation. Although it is a fast and cheap method with providing a contamination-free surface, there is a size distribution of particles and not all of them are thin enough for doing TEM observations, which makes this method unfavorable. In addition, applying this method on ductile powder particles can cause a change in their microstructure due to the induced plastic deformation. Therefore, this method can only be applied to brittle powder particles, which have an easy cleavage plane which makes the microstructure of powder particles almost unaffected in the reducing powder particles size process. Making a composite of powder particles by embedding them in a metallic matrix or an epoxy resin matrix is a more common way to prepare TEM specimens out of powder particles.
This method has been applied to a large range of powder particles by a number research groups and using a great range of metals and epoxy resins as matrices [53, 54]. 34! In 1978, Robert D. Field and Hamish L. Fraser proposed a method for preparing Ni-based superalloy powder particles by electroplating nickel (Ni) on both sides of the particle powders for 0.1 mm in thickness in a standard plating solution following with punching a 3 mm disc and electro jet-polishing the disc [55]. Later, in 1987, Kirchoff et al. also applied a TEM specimen preparation method on both aluminum powder particles and high strength aluminum powder metallurgy alloy based on depositing nickel around powder particles on a steel substrate.
Afterwards, the formed film containing powder particles embedded in a matrix of nickel was detached from the steel substrate. Although the nickel matrix can support the individual aluminum powder particles, it is a tedious and time consuming method since it requires great control over pH and temperature and also requires uniform current to form the appropriate final foil [56].
A.Montone et al. also reported a method for preparing TEM specimens of powder particles by embedding the powder particles in a metallic matrix based on the plastic follow of a soft metal sheet which can cause a cold welding between the soft metal sheets and powder particles to make a self-supporting structure by using a small hand-driven hydraulic press. This procedure followed with grind polishing and ion beam milling. They applied this method to aluminum (Al) and silver (Ag) powder particles to embed them in a copper (Cu) matrix. Although this method is very fast and easy to apply, it is applicable for powder particles with strong mechanical behavior in which the microstructure can remain unaffected under the required stress for cold welding of powder particles to matrix [53].
In 2008, Xu et al. reported a method for preparing TEM specimens out of hard ceramic powders which they applied to two hard ceramic materials, Mg10Ir19B16 and Na0.5CoO2
35! powder particles. In this method, which is based on mechanical alloying and pressing, the hard ceramic powders first mixed with soft powders of copper (Cu) and then pressed to form a pellet.
The soft copper powder can provide good encapsulation and adhesion for powder specimens.
They also point out that the ion-milling rate difference between copper and ceramic powders is much less than that of organic glues and ceramic powders. Therefore it can facilitate making thin
TEM specimens. Also, the presence of copper as a matrix with a good conductivity prevents the occurrence of any charging effect [57].
Different research groups reported methods which are based on embedding powder particles in epoxy resin or silver epoxy resin as conductive epoxy resin followed with ion-milling or FIB for powder particles sample preparation. Yang et al. prepared TEM specimens of metal particles by embedding them in Epo-Tek H20E silver epoxy followed with electrochemical polishing and a final ion-beam thinning technique [58]. Cariney et al. reported a sample preparation method based on embedding powder particles in resin, followed with focused ion- beam (FIB), applying this method for preparing a TEM specimen out of FeAl and WC powder particles [59]. Guilemany et al. also performed the same method on TiC-WC-Ni and TiC-Ti-Ni by embedding in epoxy but they used low viscosity epoxy resin following with ion beam milling
[60].
2.5! Purpose of Study
Although different research groups worked on α''-Fe16N2, as a phase with possible giant magnetization, the discrepancies among their results make the possibility of giant magnetization inconclusive and the mystery of giant magnetization of α''-Fe16N2 still remains more than 40
36! years later. As Takahashi stated in 2000, “To understand this chaotic situation of the α''-Fe16N2 problem, the most important thing is to obtain more profound knowledge of the microstructure itself.” [24] However, very few research groups did microstructural studies on iron nitrides including α''-Fe16N2 phase. To my knowledge, the microstructural studies that these few research groups have done, were mostly on nitrogen ferrite and steels, not iron nitride powders, and there is also a discrepancy among their observations themselves.
This study focused on microstructural characterization of iron-nitride powders which has been developed by Z.Feng by using Jack’s method and using nitrogen austenite as a precursor for applications in electric vehicles and renewable power generators. As a part of this project, we have applied systematic characterization methods, including transmission electron microscopy
(TEM), X-ray diffraction (XRD) for structural studies. Vibrating sample magnetometer (VSM) used for magnetic property measurements on as received sample, the quenched sample, the cryomilled sample and the heat-treated sample (see table I for details) by S.Lan. In part of phase identification and structural analysis by TEM, we have used selected area electron diffraction
(SAED), bright field (BF) imaging, dark field (DF) imaging, scanning transmission electron microscopy (STEM), and Astar for forming virtual dark field imaging on these samples. These experimental results may shed light upon variable results obtained regarding to magnetic properties and microstructural properties of α''-Fe16N2 phase by other research groups so far.
37! Chapter 3 -! Materials and Experimental Methods
3.1! Materials and Nitriding Process
The nitriding process starts with water atomized AHC 100.29 (α-Fe commercial powder), which is passed through a sieve with a diameter of around 20 µm in a designed nitriding reactor by D. Matthiesen. Prior to the start of the nitriding process, the oxide layer is removed from the surface of the starting material (AHC 100.29) by reducing it at elevated temperatures in pure hydrogen. The nitriding process is done in a horizontal reactor system which was constructed by
D. Matthiesen and is shown in Figure 12(a) and also a schematic (by D.Matthiesen) of this reactor is shown in Figure 12(b). There is a quartz boat where the starting material uniformly spread on it. The quartz boat is also connected to a thermocouple. A hydraulic piston used in order to control the place of this thermocouple can be inserted in high temperature zone or into chill-zone inside the processing tube. The gas-mixture composition is also monitored by the mass flow controllers. The system is equipped with a residual gas analyzer (RGA) to define how much gas has been decomposed. There is also a sodium filled Inconel heat pipe which is able to create an isothermal zone.
a
38! b
Figure 12: (a) Nitriding reactor, (b) Schematic of nitriding reactor by D.Matthiesen
The process schematic of producing bulk samples of α''-Fe16N2 according to Jack’s route, is shown in Figure 13. A gas mixture of NH3 and H2 has been used as the nitriding atmosphere during this process. The nitriding process takes 3600s (1h) as it has been normalized for optimum conditions. Different parameters such as, nitriding temperature, nitriding potential Kn, quench temperature and quench rate play key roles on final product composition [48]. The optimum conditions which results in forming highest fraction of nitrogen-austenite with maximum content of nitrogen interstitial atoms happens at a gas mixture of 0.11 NH3+0.89 H2 at
923.15K (650ºC), as reported by Z. Feng. In the next step, iron nitrided powder particles are quenched from 923.15K (650ºC) to a temperature below 523.15K (250ºC) in 60s (1min). As a result, γ-phase with nitrogen solid solution forms, which is also called the nitrogen-austenite phase. The resulting sample is then cryomilled at liquid nitrogen temperature for an overall time of 600s (10min) at a frequency of 30 Hz to form nitrogen martensite phase (Figure 14). The resulting sample is then heat-treated at 403.15K (130ºC) for 7200s (2h) to form ordered martensitic phase of α''-Fe16N2. All nitriding process and optimization experiments have been done by D. Matthiesen and Z. Feng. Using particle powders, can provide higher surface area which means there are higher surface area for nitrogen diffusion in nitriding process. Also, if a 39! partial transformation occurs and not all particles transformed to the final phase, there is possibility of being able to separate them by using a super magnet. Also, it’s an inexpensive method for producing α''-Fe16N2 and it offers a high chance to rapidly commercialize upon success.
α1Iron!(α1Fe) Nitridation
γ phase!with!nitrogen!solid!solution!(γ1Fe1N) Cryomill
Disordered!martensitic!phase!(α') Heat1Treatment
Ordered!Martensitic!Phase!(α''1Fe16N2)
Figure 13: Schematic of producing bulk samples of α''-Fe16N2
40!
Figure 14: Cryomilling machine
Table I: Processing conditions of samples Index Sample Sample State Nitridation Nitridation Vol% Vol% Milling Milling Tempering Tempering
Time (s) Temp (K) NH3 H2 Time (s) Frequency (Hz) Time (s) Temp (K)
1 AHC as-received ------100.29
2 AHC quenched 3600 923.15 11 89 - - - - 100.29
3 AHC cryomilled 3600 923.15 11 89 600 30 - - 100.29
4 AHC heat-treated 3600 923.15 11 89 600 30 7200 403.15 100.29
3.2! Characterization Technique
In this study, X-ray diffraction (XRD), scanning electron microscopy (SEM), vibrating sample magnetometer (VSM) and transmission electron microscopy (TEM) were used to characterize the samples.
3.2.1! X-Ray Diffraction (XRD)
41! X-ray diffraction experiments was applied on as-received sample using Scintag X-1
advanced X-ray diffractometer with a Cu Kα ("#u$%1= 0.1541 nm) source. A peak of a standard sample of alumina was acquired to make sure the machine is running properly. The XRD experiments on the quenched sample, the cryomilled sample and the heat-treated sample have
been carried out using Bruker Discover D8 X-ray diffractometer used with a Co Kα ("#o$%1=
0.1789 nm) X-ray tube.! This technique was used primarily for phase identification and measuring the lattice expansion in nitrogen-austenite phase (γ-(Fe-N)). Phase identification has been done primarily by measuring the d-spacing (dhkl) based on Bragg’s equation:
" = 2(ℎ*+,-.Θℎ*+ [49] [ 4]
λ is the wavelength of the incident X-ray and θhkl is the diffraction angle in Bragg’s equation above. For materials with a cubic structure, the lattice parameter is given by:
2 2 2 0ℎ*+ = (ℎ*+ ℎ + * + + [50] [ 5]
And for materials with a tetragonal structure, the interplanar spacing is given by:
1 ℎ2+*2 +2 = 2 + [50] [ 6] (2 02 32
The volume fraction of each phase in samples with more than one phase can also be calculated by this technique if they do not have overlaps, based on the fact that the intensity of peaks in an X-ray diffraction pattern is directly proportional to concentration of the component of phases producing it. JEMS software and crystal diffract software were used for simulating and
X-ray diffraction patterns of different phases of iron and iron-nitrides with given crystals.
42! 3.2.2! Vibrating Sample Magnetometer (VSM)
A Lakeshore Model 4710 Vibrating Sample Magnetometer was used to measure the hysteresis loops of samples in this study. Basically, this technique works on Faraday’s Law of
Induction in a way that a changing magnetic field will produce an electrical field [51]. The measured electrical field contains some information about the changing magnetic field. Also a uniform magnetic field H, will induce a magnetic moment M, in the sample, and magnetic properties of these samples can be measured as a function of magnetic field. The measured magnetic properties from the obtained hysteresis loops are saturation magnetization, coercivity , and remanent magnetization [52]. The VSM experiments have been done by S.Lan at Prof.
M.Willard’s research group in Materials Science and Engineering Department at Case Western
Reserve University.
3.2.3! Scanning Electron Microscope (SEM)
The Nova NanoLab 200 FEG-SEM/FIB, a high-resolution field emission scanning electron microscopy (SEM) with a focused ion beam (FIB) system, was used both for observing the morphology of AHC 100.29 iron powders and TEM sample preparation of samples by focused ion-beam (FIB). Cross-sectional samples of embedded iron nitride powder particles in an epoxy resin matrix and a in a commercial soldering alloy matrix, both were observed by SEM as well.
3.2.3.1! TEM Sample Preparation
Three different methods have been applied for preparing a TEM specimen out of iron- nitrided powder particles. These methods are: 1) lift out a slice of a single powder particle and thin it via FIB in a dual beam scanning electron microscope of Nova nano-lab, 2) embedding 43! iron-nitrided powder particles in Durcupan (an epoxy resin), and 3) embedding iron-nitrided particle powders in a commercial soldering alloy following with grind polishing and FIB.
To lift-out a slice of a single particle by using focused ion beam, Pt was deposited on the region of interest (Figure 15(a)) and then trenches were made at the front and back sides of the deposited Pt area. The sample were attached to an Omni-probe micromanipulator and cut from the particle and then lifted out. The slice was attached to a half copper TEM grid and thinned to become electron transparent as it can be seen in Figure 15(b).
a b
Figure 15: Lift-out protocol of a slice of a single particle
In second method, after embedding the powder particles in Durcupan (an epoxy resin) and curing it in an oven for around 8×104s at 340K-350K (70ºC- 80ºC), a slice was cut by using a precision linear saw with primary thickness of 150 µm and the surface was polished using a grinding polishing and the thickness was reduced to 46 µm and was observed under SEM. The result relating to this method is presented in chapter 4.
In third method, we applied a novel method for powder particles sample preparation based on embedding iron-nitrided powder particles in a metallic matrix. In this method which was proposed by F.Ernst, iron nitride powders particles are embedded in a commercial soldering 44! alloy via a home type solder instrument following with grind polishing the surface of the composite and FIB. The result relating to this method is presented in chapter 4.
3.2.4! Transmission Electron Microscope (TEM)
The Tecnai TF30 ST, a high-resolution analytical transmission electron microscope, operated at a high accelerating voltage of 300 kV equipped with high-angle annular dark-field (HAADF) detector for "Z-contrast" imaging, an XEDS system by EDAX with a Li-drifted Si detector with an energy resolution of 130 eV and a post-column imaging energy filter (GIF 2002 by Gatan) was used to study the microstructure of iron-nitrided samples. This microscope is also equipped with
Astar, which is mainly use as an automatic crystallographic indexing, and orientation/phase mapping tool. It consists of a CCD camera placed in front of the TEM fluorescent screen and a digital scan generator with scanning step of 0.1 nm to 100 nm in TEM, precession angle of 0°-4° and frequency of 0.1-2 kHz. Although it is mainly use as an automatic crystallographic indexing, and orientation/phase mapping tool, it is also able to create virtual dark field and bright field by using a virtual aperture and selecting particular reflections from a number of acquired electron diffraction spot patterns [53, 54].
A combination of bright field imaging (BF), dark field imaging (DF), selected-area electron diffraction (SAED), scanning transmission electron microscopy (STEM) technique and virtual
BF/DF imaging were mainly used to study the microstructure of both the matrix and the precipitations in these samples. Java electron microscope simulator (JEMS) software and Crystal
Diffract software were used for simulating and solving electron diffraction patterns of iron- nitrided TEM samples with given crystals. 45! Chapter 4 -! Results
Four samples of AHC 100.29 including as-received, quenched, cryomilled, and heat-treated
(table I), have been characterized by XRD, VSM, and TEM. VSM measurements done by S.Lan.
4.1! XRD Results
Figure 16 shows obtained XRD diffractogram of as-received sample (table I) by Z.Feng and simulated XRD diffractogram of α-iron for copper (Cu) source using crystal diffract software.
As can be seen, all peaks are matched with α-iron and no extra peak has been observed.
AHC 100.29 as received Simulated XRD for alpha-iron
022
Figure 16: XRD diffractogram of AHC 100.29 as-received and simulated diffractogram of α- iron
Figure 17 shows simulated XRD diffractograms for, γ-iron (austenite), disordered nitrogen martensite (α'-Fe8N) and ordered nitrogen martensite (α''-Fe16N2) respectively from top to
46! bottom using crystal diffract software for cobalt (Co) source and plot range (2θ) of 20º to 105º.
Simulation of crystal structures have been done by M.Willard.
022
Figure 17: simulated XRD diffractograms of γ-iron, α'-Fe8N, and α''-Fe16N2
Figure 18 shows the XRD diffractograms obtained from quenched sample, cryomilled sample and heat-treated sample at room temperature. For details on processing conditions of these sample, see table I. The first XRD pattern belongs to the quenched sample. By comparing the XRD diffractogram of this sample and the simulated XRD diffractogram of γ-iron, it can be seen that all three peaks (111), (002) and (022) have shown up in the XRD diffractogram of this sample for plot range of 15º to 105º. The positions of these peaks has been changed and there is 47! shift to the left due to the lattice expansion in the gamma phase and formation of expanded austenite. The position of these peaks in the above diffractogram are: (111) at 50.3º, (002) at
58.6º,and (022) at 87.6º. The second diffractogram belongs to cryomilled sample (table I).
Comparing the second graph with the simulated XRD diffractogram of different iron and iron- nitrided phases, it can be seen that peaks matched with the (011), (110), (002), (020), (112) and
(121) peaks of the α'-Fe8N can be seen in this pattern. Also the shifted peaks of gamma iron (γ-
Fe) exist at the same positions of previous sample with much lower intensities. The third XRD diffractogram belongs to heat-treated sample (table I). The shifted peaks of gamma iron (γ-Fe) can also be seen in this diffractogram at the same positions as the quneched and cryomilled samples. The peaks at 98.6º, 77.6º, 52.7º, 50.3º appeared at the same positions as can be seen in the XRD diffractogram of the cryomilled sample. However, two peaks in the XRD diffractogram of the heat-treated sample have shifted to the left in comparison to the peaks in the cryomilled sample.
48! Quenched sample Cryomilled sample
(111)γ Heat-treated sample
(002)γ
(022)γ
(111)γ (011)α' (110)α'
(121)α' (002)γ (002)α' (020)α' (112)α' (022)γ
.
Figure 18: XRD patterns of quenched sample, cryomilled sample and heat-treated sample
Figure 19 shows the superimposed XRD diffractograms of the cryomilled sample and heat- treated sample . As it can be seen from the superimposed XRD diffractograms of these two samples in Figure 19, the two peaks, which appeared at 70.1 º and 93.1 º in the cryomilled sample, shifted to 69.5ºand 92.5º respectively in the heat-treated sample. The peak at 69.5º can belong to (040) of α''-Fe16N2 phase or (020) peak of α'-Fe8N at higher tetragonality (c/a) ratio and peak at 92.5º can also belong to (224) peak of α''-Fe16N2 or (112) peak of α'-Fe8N with higher tetragonality (c/a) ratio. Since there is an overlap between XRD peaks of expanded
49! austenite (γ-Fe) and nitrogen martensite (α'-Fe8N) with ordered nitrgoen martensite (α''-Fe16N2), other techniques such as TEM and VSM are used to confirm the formation of α''-Fe16N2.
Cryomilled sample Heat-treated sample
(112)α' (020)α'
(040)α'' (224)α''
Figure 19: Superimposed XRD diffractograms of cryomilled sample and heat-treated sample
4.2! VSM Results
The magnetic properties of as-received sample, quenched sample, cryomilled sample, and heat- treated sample have been measured using VSM by S.Lan. Figure 20 shows the mass magnetization versus magnetic field strength (M vs. H) curves of these four samples. Table II, shows the VSM measurements including mass magnetization at 1.5×107/4π A/m (15 kOe), coercivity (A/m) and remanent magnetization (A.m2/kg) for these four samples. From Table II, it can be seen that the magnetization at 1.5×107/4π (A/m) (15 kOe) for as-received sample containing only α-iron is 212.7 Am2/kg. After nitridation, for the quenched sample this value
50! decreases to 1.9 A.m2/kg. For the cryomilled sample, this value increases to 158.4 A.m2/kg and there is a further increase to 202.5 A.m2/kg for the heat-treated sample. Coercivity also shows a great increase for the heat-treated sample in comparison to as-received sample from
5.0×103/4π Α/m to 165.1×103/4π Α/m for the heat-treated sample. Also, the remanent magnetization shows a great increase from 0.32 A.m2/kg to 8.49 A.m2/kg. The standard error in
2 these measurements is ±0.2 A.m /kg. These measurements were done by S. Lan and M. Willard.
a d
c
b
As-received sample Quenched sample Cryomilled sample Heat-treated sample
Figure 20: VSM curves of four samples of (a) As received sample (b) Quenched sample (c) Cryomilled sample (d) Heat-treated sample
51! Table II: Magnetic properties measurement of the samples Sample state Magnetization (A.m2/kg) Coercivity (A/m) Remanent Magnetization
at 1.5 ×107/4π A/m (A.m2/kg)
as received 212.7 5.0 × 103/4π 0.32
quenched 1.9 313.1 × 103/4π 0.39
cryomilled 158.4 110.4 × 103/4π 5.01
heat-treated 202.5 165.1 × 103/4π 8.49
4.3! TEM Sample Preparation Results
Figure 21 shows a SEM image of cross-sectional embedded powder particles in the epoxy resin (Durcupan) with composite thickness of 46 µm. As we can see there are some cracks which have been formed around the powder particles and therefore, powder particles do not have enough adhesion to the matrix and most probably the matrix is not able to support the powder particles during thinning.
52!
Figure 21: SEM image of cross-sectional embedded particle powders in the epoxy resin (Durcupan)
Figure 22 shows a SEM image of cross-sectional iron-nitrided powders embedded in soldering alloy. Figure 23 shows a XEDS map for iron (Fe), tin (Sn) ,and lead (Pb). As can be seen, the darker areas are related to iron-nitrided particle powders as is expected, and the powder particles have a good adhesion to the surface.
53!
Figure 22: SEM image of cross-sectional embedded particle powders in commercial solder
Figure 23: XEDS map of cross-sectional embedded particle powders in commercial solder
Figure 24(a) shows a low-mag HAADF STEM image acquired from the final TEM foil prepared with this method. As can be seen, the resulting TEM specimen has a self-supporting structure.
Figure 24(b) shows a magnified image of the areas in which XEDS scans are taken. As can be seen from the XEDS maps (Figure 25), there is no sign of redisposition of soldering alloy on iron-nitrided particles (Figure 25 (a)). The copper (Cu) peaks which can be seen in all three scans mostly come from the copper grid.
54! a
b
1 2
3
Figure 24: STEM image of cross-sectional embedded particle powders in commercial solder (a) low magnification (b) higher magnification of marked area
a
55! b
c
Figure 25: XEDS map of cross-sectional embedded particle powders in commercial solder
4.4! TEM Results
4.4.1! TEM observation of quenched sample
Figure 26(a) shows a low magnification STEM image acquired at camera length (CL) of 120 mm of TEM foil prepared from a powder particle of the quenched sample (table I) by FIB.
Because the STEM images are formed by high angle incoherently scattered electrons, they are very sensitive to the variations in atomic number of existing elements in sample (Z-contrast images); therefore, we expect the elements with higher atomic number appear brighter in comparison to areas consisting of lighter elements. For thicker parts of specimen (higher t), we also expect the electrons scatter more strongly, therefore a higher number of them are collected 56! by HAADF detector and they appear brighter in comparison to thinner area of the specimen.
Figure 26(b) and Figure 26(c) show diffraction contrast STEM images acquired from the big
(central) needle, taken from the bottom part and the tip part of the needle respectively, at camera length (CL) of 250 mm. Three needle-shaped features can be seen in this image, shown with arrows on Figure 26(a). Varying camera length in STEM mode corresponds to varying the size of objective aperture in TEM mode. Therefore, the collection angle of detector changes by changing the camera length (L) in STEM mode. Diffraction contrast arise from Bragg scattered electrons and STEM detector collect several Bragg scattered beam which reduce the effect of diffraction contrast in STEM. However, with increasing camera length (L) in STEM mode, we expect the diffraction contrast also contributed in the STEM images. Since Figure 26(b) and
Figure 26(c) were acquired at camera length (CL) of 250 mm, we expect diffraction contrast also contributed in these images. As can be seen from Figure 26(c), a contrast can be seen in big needle itself by increasing the camera length to 250 mm, which can be due to grains with different orientations or two different phases. As can be seen in Figure 26(b), the brighter areas are closer to edge of the needle rather than center of the needle.
57!
Figure 26: (a) low mag STEM image of nitride sample at CL: 120 mm (b) magnified STEM image from the bottom of main needle at CL: 250 mm (c) magnified STEM image from the tip of main (central) needle at CL: 250 mm
Figure 27 shows three bright field images taken from the bottom, middle and tip of the area containing the big (central) needle in low mag STEM image (Figure 26(a)). The contrast in the needle and around the needle could possibly be due to the strain field around the precipitates 58! which has been formed as a result of elastic deformations of the matrix mainly caused by the presence of precipitates [55]. Bend contours as an artifact of FIB sample preparation can also be seen in the matrix area of this image.
1
2
3
Top surface
Figure 27: Bright field images from bottom, middle and tip of the area containing the big needle in low mag STEM image
Figure 28(a) shows the electron diffraction pattern obtained from the marked area 1 on
Figure 27, corresponding to the bottom part of the needle. Figure 28(b) shows adjusted grey level 59! of the central area of Figure 28 (a), the obtained electron diffraction pattern. As can be seen, the superlattice reflections are lit up in this electron diffraction pattern (Figure 28(b)). Which may have caused by long-range ordering and indicating the formation of an ordered phase. However, the very low intensity of superlattice reflections is related to the difference in atomic scattering factors of the ordering elements which is nitrogen in this case and iron atoms [49]. Figure 28(c) shows the simulated electron diffraction pattern for [111] zone axis of α''-Fe16N2 crystal by using the crystal diffract software package. Comparing d-spacing (inverse of measured distances (1/d) on electron diffraction pattern) and angles from the obtained electron diffraction pattern and the simulated electron diffraction pattern confirms that this diffraction pattern corresponds to α''-
Fe16N2 phase at [111] zone axis. Figure 28(d) shows the solved electron diffraction pattern for
α''-Fe16N2 phase at [111] zone axis by using JEMS software. Diffuse spots in Figure 28(a), can also be because of local lattice distortion due to the insertion of nitrogen atoms in them.
Although, superlattice reflections lit up in Figure 28(a), this area may contain both disorder nitrogen martensite (α') and ordered nitrogen martensite. However, because of the very small misfit, we don’t see any separation among spots in Figure 28(a).
Figure 28(e) shows the obtained electron diffraction pattern from the top part of the needle.
Figure 28(f) shows the grey level adjusted of the central part of Figure 28(e). As can be seen, no superlattice reflection is seen in the obtained diffraction pattern of the top part of the needle which corresponds to a disordered martensitic phase α'-Fe8N phase at [111] zone axis. Extra marked spots (marked with black circles), could possibly indicating the small orientation difference (around 5°) between adjacent grains.
60!
a c 4.74 -1
nm 4.94 nm - 1
b
d
61! e
5°
000
f
Figure 28: (a) electron diffraction pattern corresponding to bottom part of the needle (area marked 1) (b) shows the grey level adjusted of central area (c) shows the simulated electron diffraction pattern for [111] zone axis of α''-Fe16N2 (d) solved electron diffraction pattern for α''- Fe16N2 phase at [111] zone axis by using JEMS software (e) electron diffraction pattern corresponding to top part of the needle (f) grey level adjusted of central area of part (e)
Figure 29(a) shows an obtained diffraction pattern from the matrix near the big needle (area 2 on Figure 27). Figure 29(b) shows simulated diffraction pattern for [110] zone axis of γ-Fe phase without nitrogen. Comparing d-spacing (inverse of measured distances (1/d) on electron diffraction pattern) and angles from obtained electron diffraction pattern and simulated electron diffraction pattern confirms that this diffraction corresponds to the expanded γ-iron phase at the
[110] zone axis. Figure 29 (c) shows the solved electron diffraction pattern for [110] zone axis of
γ-Fe by using JEMS software. The lattice expansion calculated from obtained diffraction pattern is estimated 2.3%. Marked extra spots which are also available around the transmitted beam, can be interpreted as a result of double diffraction due to sample thickness.
62! a b
54.7o
5.62 nm-1
c
Figure 29: (a) shows an obtained diffraction pattern from area 2 marked on Figure 27(b) shows the simulated diffraction pattern for the [110] zone axis of γ-Fe phase (c) solved electron diffraction pattern for the [110] zone axis of γ-Fe phase by using JEMS software
63! DF image acquired from one of the superlattice reflections to see whether the extra spots arise from ordering or not. Figure 30 shows the dark field image obtained from one of the superlattice reflections as marked on diffraction pattern, belonging to {112} planes.
Corresponding areas to this spot are areas mostly close to the edge of the needle which are lit up in the obtained dark field image. These areas are very close to the brigth area in the STEM image
(Figure 26(b)).
Figure 30: Dark field image obtained from one of the super-lattice reflections belonging to {112} planes
Figure 31(a) shows a bright field image of the top area of the quenched sample containing two small needle-shaped features. This area corresponds to top-right part in low magnification
STEM image in Figure 26(a). Figure 31(b) shows the electron diffraction pattern obtained from
64! the surrounding matrix which is marked on the image as area 1. Figure 31(c) shows the electron diffraction pattern obtained from the area 2, containing both the matrix and smaller needle.
Figure 31(d) shows the simulated electron diffraction pattern of α'-Fe8N for [100] zone axis. As can be seen, the electron diffraction pattern of the surrounding matrix corresponds to expanded austenite for [110] zone axis comparing to the simulated diffraction pattern for the [110] zone axis of γ-Fe phase (Figure 29(b)) with the estimated lattice expansion of 2.3%. As can be seen from Figure 31(c), extra spots which are shown with white circles correspond to the disordered nitrogen martensitic phase at the [100] zone axis comparing with the simulated electron diffraction pattern of α'-Fe8N for [100] zone axis in Figure 31(d). Figure 31(c) have revealed
that, there is a possible preferred orientation relationship of [100]α'" [101]γ and (011)α'" (111)γ between the martensitic lath and expanded austenite matrix in this sample which shows closed- packed planes in each phase should be parallel to each other and it can possibly indicate the martensitic laths are coherent with the expanded austenite matrix and this is matched with a
Nishiyama-Wasserman (N-W) orientation relationship [56].
a
1
2
65! b
c
d
84.42o
Figure 31: (a) Bright field image of top right part of sample including smaller needles and matrix (b) electron diffraction pattern of matrix (c) electron diffraction pattern of marked area 66! containing small needle and surrounded matrix (d) simulated electron diffraction pattern of α'- Fe8N for [100] zone axis
4.4.2! TEM observations of cryomilled sample
Figure 32(a) and Figure 32(b) show STEM images obtained at camera length (CL) of 250 mm from cryomilled sample (table I) in order to include the diffraction contrast which can come up from two different phases or two different orientations in addition to Z-contrast. As can be seen, the microstructure is mainly composed of groups of more or less parallel martensitic laths.
Dark spots can be due to beam damage during image acquisition in Figure 32(a). Figure 32(c) shows the electron diffraction pattern of the marked area by a black circle on Figure 32(b) on the curved lath. Comparing Figure 32(c) with Figure 32(d) which is the simulated diffraction pattern of austenite (γ-Fe) without nitrogen for [110] zone axis, confirms that this diffraction pattern corresponds to an area containing expanded austenite phase (γ-Fe) with 2.3% of lattice expansion. Extra reflections appeared at the half reciprocal spacing of expanded austenite (200) reflections, resulting from the double diffraction, which can be due to the point that the sample on that area is not thin enough (it is also brighter in STEM image). As can be seen, the retained austenite phase also formed lath shaped features in this sample. Areas shown with arrows on
Figure 32 (c) can possibly indicate areas with high density of defects. Figure 32(e) shows the bright field image corresponding to area shown on STEM image in Figure 32(a).
67! a
b
68! c d
54.7o
5.49 nm-1 5.62 nm-1
e
Figure 32: (a) First and (b) second STEM images of cryomilled sample at CL:250 mm (c) simulated electron diffraction pattern of γ-iron for [110] zone axis (d) electron diffraction pattern of marked area in b (e) bright field image of corresponding area in part (a)
69! Figure 33(a) shows a TEM image of the cryomilled sample and the area that the electron diffraction pattern in Figure 33(c) obtained. Figure 33(b) shows the simulated diffraction pattern of α'-Fe8N for the [113] zone axis. Comparing d-spacing (inverse of measured distances (1/d) on electron diffraction pattern) and angles from obtained electron diffraction pattern and simulated electron diffraction pattern can confirm that this electron diffraction pattern match with the simulated electron diffraction pattern of α'-Fe8N phase for the [113] zone axis. Figure 33(d) shows the solved electron diffraction pattern for [113] zone axis of α'-Fe8N by using JEMS software. Spots splitting or satellite which is marked by black circle on Figure 33 (c) can occur because of the presence of planar defects such as dislocations or regular arrays of crystal defects
[50].
a
70! b
c
d
Figure 33: (a) martensitic lath in cryomilled sample (b) electron diffraction pattern from marked area (c) simulated diffraction pattern of α'-Fe8N for the [113] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [113] zone axis by using JEMS software
Figure 34(a) shows the area that the electron diffraction pattern obtained (area 1). Figure
34(b) shows obtained electron diffraction pattern of area marked as area 1. Figure 34(c) shows the simulated diffraction pattern of α'-Fe8N for the [111] zone axis. Comparing d-spacing
(inverse of measured distances (1/d) on electron diffraction pattern) and angles from obtained electron diffraction pattern and simulated electron diffraction pattern can confirm that this diffraction pattern matches with electron diffraction pattern of α'-Fe8N phase for the [111] zone axis. Figure 34(d) shows solved electron diffraction pattern of α'-Fe8N for [111] zone axis by 71! using JEMS software. As can be seen, the experimental electron diffraction pattern (Figure
34(d)) shows a match with simulated diffraction pattern for α'-Fe8N phase at the [111] zone axis at low index planes. However, it shows a slight derivation to smaller d-spacing for higher index planes as shown by an arrow in Figure 34(d) which can probably be due the nitrogen deficient regions and higher concentration of nitrogen atoms in lower indexed planes. Diffuse spots in
Figure 34(b), can also be because of local lattice distortion due to the insertion of nitrogen atoms in them [55].
a
1
72! b c
58.5o
58.42o
1 - 5.00 nm 5.00
d
Figure 34: (a) shows a TEM image containing the area the electron diffraction pattern obtained (b) obtained electron diffraction pattern (c) simulated diffraction pattern of α'-Fe8N for [111] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [111] zone axis by using JEMS software
Figure 35(a) shows the martensitic lath and also thin lath-shaped feature from which the electron diffraction pattern obtained. Features with the same appearance are also marked with arrows in Figure 32(e). Figure 35(b) shows the obtained electron diffraction pattern from the
73! marked area. Figure 35(c) shows the simulated diffraction pattern of α-iron for the [100] zone axis. Comparing d-spacing (inverse of measured distances (1/d) on electron diffraction pattern) and angles from obtained electron diffraction pattern and simulated electron diffraction pattern can confirm that this diffraction does not show any appreciable tetragonality and it is very close to the electron diffraction pattern of the α-Fe phase for the [100] zone axis. Figure 35(b) shows extra spots appeared at half distance of (110) planes corresponding to the (110) planes of α''-
Fe16N2 which can probably be due to the formation of very small fraction of α''-Fe16N2. As can be seen in Figure 35(b), some spots are stretched into arches, which can possibly interpret as grain rotation introduced by interstitial nitrogen atoms and as a result of deformation of austenitic matrix during martensitic transformation [57]. Figure 35(d) shows solved electron diffraction pattern of α-Fe for [100] zone axis by using JEMS software.
a
74! b c
d
Figure 35: (a) small laths and the area that the electron diffraction pattern obtained (b) electron diffraction pattern of marked area (c) simulated diffraction pattern of α-iron for [100] zone axis (d) solved electron diffraction pattern of α-iron for [100] zone axis by using JEMS software
4.4.3! TEM observations heat-treated sample
Figure 36(a) and Figure 36(b) show diffraction contrast STEM images of heat-treated sample
(table I) acquired at camera length (CL) of 120 mm. Figure 36(c) shows HAADF-STEM image
75! of the same area but in higher magnification which has been recorded at CL of 80 mm. By increasing the camera length from 80 mm to 120 mm, diffraction contrast also contributed to make the STEM images. As can be seen, needles (Figure 36(b)) show an internal contrast appearing darker in the central parts and brighter in the areas near the surface of these needles which can ascribing as different orientation or different phases. The contrast between central part and surface areas of these needles increased by increasing the camera length from 80 mm to 120 mm can be due to different orientation or different phases. As can be seen, this series of needles appeared more or less parallel. Dark spots appeared in STEM images can be due to beam damage during image acquisition.
a c b
Figure 36: STEM images of heat-treated sample (a) and (b) obtained at CL: 120 mm (c) obtained at CL: 80 mm
Figure 37 (a) (b) (c) shows BF images of three sets of needle shaped precipitations in the heat-treated sample acquired by inserting the objective aperture on back focal plane around the 76! direct beam and contribute in making image. Generally, the contrast appears in BF images are due to mass-thickness for not thin enough samples and diffraction contrast. As it can be seen, three sets of needle-shaped and lath-shaped features appeared after heat-treating the sample. It can be seen in Figure 37(d) that one set of lath-shaped features also composed of smaller size laths in themselves. These three sets can have different composition or they might have same structure only viewed in different orientation. If we categorize these needles to three different sets, each set is parallel to each other.
77! a b Set 2
Set 1
Set 3
`
c d
Figure 37: (a) (b) (c) (d) bright field (BF) images of three sets of needles in tempered sample in different magnifications
Figure 38(a) shows a TEM image of the sample after tilting. Figure 38(b) shows the
diffraction pattern obtained from the marked area in Figure 38(a). Figure 38(c) show the
simulated diffraction pattern of α'-Fe8N for the [111] zone axis. Comparing the angles and d-
78! spacing of obtained diffraction pattern with simulated diffraction pattern of α'-Fe8N can confirm that this diffraction pattern is also corresponding to α'-Fe8N in the [111] zone axis. Figure 38(d) shows solved electron diffraction pattern of α'-Fe8N for [111] zone axis by using JEMS software.
The streaks which are shown by black circles can be due to the shape effects of precipitations in this sample indicating a high density of crystallorgraphically orientated defects such as precipitates in this case [50].
a a a
79! b c 5.00 nm - 1
d
Figure 38: (a) TEM image after tilting (b) electron diffraction pattern of marked area (c) simulated electron diffraction pattern of α'-Fe8N for [111] zone axis (d) solved electron diffraction pattern of α'-Fe8N for [111] zone axis by using JEMS software
80! Figure 39(a) shows a TEM image after tilting the sample. Figure 39(b) shows an electron diffraction pattern acquired from the marked area on Figure 39(a). Figure 39(c) shows the adjusted grey level of the acquired electron diffraction pattern. As it can be seen the superlattice reflections appeared clearly in Figure 39(c) which can ascribed as existing a long range ordering in the SADP. Figure 39(d) shows the simulated diffraction pattern for the [100] zone axis of α''-
Fe16N2. As can be seen, the superlattice reflections in the obtained electron diffraction pattern are revealed at predicted positions from simulated diffraction pattern for the [100] zone of α''-
Fe16N2. Also, comparing d-spacing (inverse of measured distances (1/d) on electron diffraction pattern) and angles from obtained electron diffraction pattern and simulated electron diffraction pattern can confirm that this electron diffraction pattern corresponds to α''-Fe16N2 phase at [100] zone axis. Figure 39 (e) shows solved electron diffraction pattern of α''-Fe16N2 for [100] zone axis by using JEMS software. The streaks which can be seen on almost all main spots in the same direction (Figure 39b) and shown by black circle, can be due to the shape effects of precipitations in this sample indicating a high density of crystallorgraphically orientated defects such as precipitates in this case or the lattice strain associated with them [50].
81! a
b d
o 95.43 c
82! e
Figure 39: (a) TEM image of tilted sample, marked area shows the corresponding area of obtained diffraction pattern (b) electron diffraction pattern of marked area (c) grey level adjusted electron diffraction pattern to reveal superlattice reflections (d) simulated electron diffraction pattern of α''-Fe16N2 for [100] zone axis (e) solved electron diffraction pattern of α''-Fe16N2 for [100] zone axis zone axis by using JEMS software
Figure 40 shows dark field images obtained from three superlattice reflections in obtained electron diffraction pattern corresponding to α''-Fe16N2 at [100] zone axis to see whether the extra spots arise from ordering or not . These spots are corresponding to {011} planes in Figure
40(a) and Figure 40(b). Figure 40(c) corresponds to {002} planes. As can be seen in Figure
40(a), areas near the surface of primary needle light up. These areas correspond to brighter areas in the STEM image (Figure 36(b)). Also, some parts of the matrix became bright as well. The central part of these needles remained dark and the size of central black part is more or less the same size of the central part of other sets which initially named them as three sets. The parts of the matrix which also lit up in Figure 40(a) correspond to brighter areas in STEM image (Figure
36(a)).
83!
a
b c
Figure 40: Dark field images corresponding to (a) and (b) {011} planes (c) {002} planes in α''- Fe16N2 electron diffraction pattern at [100] zone axis
84! Figure 41(a) shows a TEM image of sample after tilting. Figure 41(b) shows the electron diffraction pattern of the area marked on the Figure 41(a). Figure 41(c) shows a simulated diffraction pattern of α''-Fe16N2 phase for the [103] zone axis. Comparing obtained diffraction pattern with simulated diffraction pattern for the [103] zone axis of α''-Fe16N2 confirms that the obtained diffraction pattern is also corresponding to α''-Fe16N2 phase for the [103] zone axis.
Figure 41(d) shows solved electron diffraction pattern of α''-Fe16N2 for [103] zone axis by using
JEMS software. The extra spots marked by black circle in Figure 41 (d) can possibly coming from the matrix.
a
85!
d
Figure 41: (a) area of obtained electron diffraction pattern (b) corresponding electron diffraction pattern of marked area (c) simulated electron diffraction pattern of α''-Fe16N2 for the [103] zone axis (d) solved electron diffraction pattern of α''-Fe16N2 for [103] zone axis by using JEMS software 86! Figure 42 shows a dark field image corresponding to the {301} planes as marked on the obtained electron diffraction pattern corresponding to ordered nitrogen martensite (α''-Fe16N2) for the [103] zone axis. As can be seen, some areas in matrix and a set of needles are mainly what light up.
Figure 42: dark field image corresponding to spot from the {301} planes in α''-Fe16N2 electron diffraction pattern at the [103] zone axis
Figure 43 (a) (b) show virtual TEM images and corresponding diffraction patterns of area near to the surface. Figure 43 (c) (d) show virtual TEM images and corresponding diffraction patterns of central areas of the needles. Figure 43 (e) shows virtual TEM images and corresponding diffraction patterns of bright areas of the matrix. As can be seen, superlattice reflections indicating a possible long range ordering can be seen in DPs of area closer to the surface of the needles and also the bright area of matrix.
87! c a b d
e
Figure 43: (a) (b) virtual TEM images and DPs of area near to the surface of needle (c) (d) virtual TEM images and DPS of central area of needles (e) TEM image and DP of bright area of matrix
Figure 44 (a) (b) (c) shows virtual dark field images acquired by using Astar from marked superlattice reflections in obtained DP (Figure 44 (e)). Figure 44 (d) shows virtual dark field image acquired from marked main spot on DP by using Astar. As can be seen, by selecting the superlattice reflections which can possibly indicate a long range ordering, area near to the surface of the needle and some areas of matrix lit up. As can be seen in Figure 44 (d) which has been acquired by selecting the marked main spot, the central area of the needles become bright.
This observation in virtual dark field images obtained by selecting superlattice reflections are consistent with our observation by acquiring dark field images from superlattice reflections in real diffraction patterns.
88! a b c
d
Figure 44: (a) (b) virtual dark field images corresponding to superlattice reflections marked on diffraction pattern (d), (c) virtual dark field image corresponding to marked main spot on diffraction pattern
89! Chapter 5 -! Discussion
5.1! XRD Results
XRD result of the nitrided sample in Figure 18 shows the existence of a single phase of austenite(γ-Fe) in this sample while a significant peak shift of austenite to lower diffraction angles, indicating lattice expansions of austenite during nitridation and forming nitrogen austenite has been observed. No extra peak indicating an existence of other phases was observed in XRD results of this sample. The lattice parameter of expanded nitrogen austenite in this sample was measured by using crystal diffract software by comparing the peak positions of (111) at 50º, (002) at 58.6º and (022) peaks at 87.8º in the obtained result and peak positions in simulated X-ray diffraction pattern for austenite (γ-Fe). The measured value for the lattice parameter of nitrogen austenite in this sample was 0.3649 nm. The nitrogen content for nitrogen austenite in this sample, calculated by using the Wriedt equation [9] and using 0.3572 nm for the lattice parameter of elemental austenite γ-Fe without nitrogen, is (9.8 ± 0.4)at.%, less than 10.3 at% which is the maximum nitrogen content for nitrogen austenite at 923.15K (650°C) according to Fe-N phase diagram.
The second graph in Figure 18, which represents the XRD result of the cryomilled sample, shows this sample was composed of nitrogen austenite (γ-(Fe-N)) as retained austenite from its previous step and nitrogen martensite (α'), which has been formed during cryomilling at liquid nitrogen temperature. The tetragonality (c/a ratio) of formed nitrogen martensite in this sample was measured by crystal diffracts software as 1.09. The retained austenite phase as expected has the same lattice parameter of its parent austenite in the quenched sample, because the peak positions of retained austenite did not change in comparison to the nitrogen austenite phase, 90! which was formed in the previous step. Therefore, the retained austenite has the same nitrogen content as its parent austenite of the previous step. The volume fraction of retained austenite, according to the XRD result of this sample is around 9%, although it was expected that all of the nitrogen austenite phase would have been transformed into nitrogen martensite during cryomilling at liquid nitrogen temperature. There could be several reasons for the observation that there is a small amount of retained austenite in cryomilled sample. One such reason is that by increasing the nitrogen content, the Ms (martensitic start temperature) will decrease.
Moreover, the under-cooling was not sufficient to transform all of the nitrogen austenite phase to disordered nitrogen martensite phase. However, 11.1 at.% of nitrogen should be primarily inserted into nitrogen austenite and we can get 100 % of α''-Fe16N2 after heat-treatment. The maximum nitrogen content for the quenched sample is (9.8 ± 0.4) at.%, so it is not expected that all nitrogen martensite would transform into α''-Fe16N2.
The third graph in Figure 18 corresponds to the XRD diffractogram of the heat-treated sample. Although it is expected to get ordered nitrgoen martensite (α''-Fe16N2) in the heat-treated sample, traces of retained austenite and disordered nitrogen martensite could be seen at the same positions, as it was observed in the quenched sample and the cryomilled sample. There is an overlap between peaks of ordered martensite and nitrogen martensite and the expanded gamma phase. In Figure 19, which shows the superimposition of two XRD graphs of the cryomilled sample and the heat-treated sample, it can be seen that, in addition to retained austenite peaks, some other peaks appeared at the same positions of the previous sample (the cryomilled sample) and two peaks have shifted to the left. Peaks at 98.6º, 77.6º, 52.7º, 50.3º appeared at the same positions of previous sample (cryomilled sample). Two peaks appearing at 70.1º and 93.1º in the cryomilled sample shifted to 69.5º and 92.5º respectively in the heat-treated sample. The peak at 91! 69.5º possibly can belong to the (040) peak of α''-Fe16N2 phase or the (020) peak of α'-Fe8N at higher tetragonality (c/a) ratio and the peak at 92.5º can also belong to the (224) peak of α''-
Fe16N2 or the (112) peak of α' with higher tetragonality (c/a) ratio compared to the α' phase formed in the cryomilled sample, and there is an overlap between XRD peaks of the expanded austenite (γ-Fe) and nitrogen martensite (α'-Fe8N) with ordered nitrgoen martensite (α''-Fe16N2).
These two peaks shifted to the left, which can possibly mean that by doing heat-treatment, the tetragonality increased. The tetragonality of the heat-treated sample is estimated as 1.10, according to shifts in position of these two peaks after heat-treatment. However, this is the expected value for tetragonality of α''-Fe16N2 phase. According to Jack’s report in 2000, there is a hypothesis describing that there is a pre-precipitation process before nitrogen atoms get long- range ordered in α''-Fe16N2,in which the disordered nitrogen martensite clusters together and the tetragonality increases until it reaches the tetragonality of the ordered nitrogen martensite phase and the nitrogen gets ordered and we have an α''-Fe16N2 phase [14]. Therefore, from acquired
XRD diffractograms, it can be concluded that, after heat-treatment tetragonality increased to a value corresponding to the tetragonality of the α''-Fe16N2 phase but it cannot be concluded whether nitrogen atoms get ordered in a perfect manner to form the ordered martensitic phase of
α''-Fe16N2 or not. For further investigation to confirm whether this phase has been formed or not, the electron diffraction patterns of TEM foil of this sample for more structural analysis will be presented.
5.2! VSM Results
92! As can be seen from Figure 20, which shows the mass magnetization versus magnetic field strength (M vs. H) curves of these samples at mass magnetization of 1.5 × 107/4π (A/m) (15 kOe), the mass magnetization at 1.5 ×107/4π (A/m) (15 kOe) decreases from 212.7 Am2/kg for primary α-iron to 1.9 A.m2/kg for quenched sample due to the formation of nitrogen-austenite.
This result also confirms the XRD data for this sample, which only shows the peaks related to expanded austenite. Also, VSM data for this sample shows that it contains almost entirely nitrogen austenite except for a very small fraction of a magnetic phase. The fraction of a second magnetic phase in this sample should be smaller than the detection limit of the XRD method, as it has not appeared in the XRD results for this sample (less than 5%), but mass magnetization for the quenched sample (1.9 A.m2/kg) at 1.5 × 107/4π (A/m) (15 kOe) shows the existence of a second phase. For the cryomilled sample, the mass magnetization at 1.5 ×107/4π (A/m) (15 kOe) increases to 158.4 A.m2/kg, which is due to the transformation of a significant portion of nitrogen austenite into nitrogen martensite. As the XRD result for this sample shows peaks related to the retained austenite, it can be concluded that nitrogen austenite did not fully transform into nitrogen martensite after cryomilling. For the heat-treated samples, the mass magnetization at 1.5 × 107/4π (A/m) (15 kOe) increased from 158.4 A.m2/kg for the cryomilled sample to 202.5 A.m2/kg for the heat-treated sample, which can be due to the formation of ordered nitrogen martensite (α''-Fe16N2) from disordered nitrogen martensite (α'). Although the
VSM can possibly confirm the formation of ordered nitrogen martensitic (α''-Fe16N2) with possible giant magnetization in the heat-treated sample, the mass magnetization at 1.5 × 107/4π
(A/m) (15 kOe) of final heat-treated sample is lower than the mass magnetization at 1.5 × 107/4π
(A/m) (15 kOe) for the primary α-Fe sample which is due to the existence of retained austenite to around 9 at.% in this sample. Coercivity also shows a great increase after heat-treating the 93! sample in comparison with the primary α-Fe powder from 5.0 × 103/4π Α/m to 110.4 ×
103/4π for the cryomilled sample and to 165.1 ×103/4π Α/m for the heat-treated sample. This increase in coercivity is a desirable property for a promising permanent magnetic material, making the heat-treated sample containing the α''-Fe16N2 phase a great candidate, as it indicates that it is stable against demagnetization. In addition, the remanent magnetization shows a increase from 0.32 A.m2/kg for the primary α-Fe sample to 5.01 A.m2/kg for the cryomilled sample after formation of the disordered nitrogen martensite and finally to 8.49 A.m2/kg for the final heat-treated sample which shows the remained magnetization after removal of the external magnetic field is higher for heat-treated sample in comparison with primary α-Fe sample.
5.3! TEM Sample Preparation Results
The powder particles are embedded in epoxy because the ion-milling rate between particles and the epoxy resin as a matrix is quite different. Therefore, there is a chance that particles are lost during the TEM sample thinning process. As is shown in Figure 21, indicating that there is not a good adhesion between powder particles and matrix. In addition, the epoxy resin is not very stable under the electron beam and there is a chance of drifting during the observation. This method probably can work for biological samples but not for metallic powder particles.
By embedding the powder particles in commercial soldering alloy followed by FIB, because the matrix has a good thermal and electrical conductivity; therefore, we expect the resulting sample to be stable under the electron beam. In addition, this method is fast and cheap in comparison to other embedding methods for preparing TEM specimens from powder particles such as electroplating. As Xu et al. stated, preparing composite samples of powder particles can
94! provide more accurate statistical information, since they can provide a bigger area for TEM observations in comparison to preparing an FIB sample of a single powder particle, which can only provide information of a small area of a single particle [58]. Embedding the powder particles in commercial soldering alloy followed by FIB can be used for preparing TEM samples of a great range of materials, except for hard ceramic powders, due to the huge difference in the ion-milling rate of hard ceramic powder particles, while this ion-milling rate difference is quite acceptable between metallic powders and soldering alloy and can provide enough adhesion between the embedded powder particles and matrix. Moreover, the XEDS spectra of this sample can reveal that we do not have any redeposition problem in this method.
5.4! TEM Results
5.4.1! Quenched sample
Low magnification STEM image of quenched sample (Figure 26) shows the formation of a few needle-shaped precipitations in this sample. STEM images at higher camera length (CL: 250 mm) showed a contrast within the central (big) needle which can be due to two different orientations or two different phases. The electron diffraction pattern taken from the bottom of the needle at the [111] zone axis revealed that the obtained diffraction pattern can be matched in angles and d-spacing with simulated diffraction pattern of α''- Fe16N2 phase for the [111] zone axis. Additionally, the superlattice reflections light up in the obtained diffraction pattern with lower intensity in comparison to the main spots. However, the electron diffraction pattern obtained from the top of this needle did not show any superlattice reflections and it can be matched with α'-Fe8N for [111] zone axis. The dark field image, which was taken from one of the superlattice reflections corresponding to {112} planes shows that some small regions
95! primarily close to the edge of the needle light up. This observation can possibly ascribe that nitrogen atoms in small regions of this needle became ordered to some extent. It seems those areas which are brighter in low magnification STEM image (Figure 26) are matched with parts of the needle that light up in the dark field image (Figure 30). These needles can be stress- induced martensite or related areas with low nitrogen content, which would imply that Ms
(martensite starting temperature) had an increase; therefore, after quenching, a very small volume fraction of nitrogen martensite had been formed. But the reason that nitrogen atoms got ordered to an extent in nitrogen martensite can be due to the fact that in these lath martensite, because of the stress fields around individual dislocation and cell walls, certain interstice positions near these defects can provide lower-energy sites for interstitial atoms such as nitrogen rather than the normal interstitial lattice positions and a local rearrangement of nitrogen atoms near to these positions occurring. That is probably why nitrogen atoms in small regions near to the surface of the main needle became ordered [59, 60].
The electron diffraction pattern which has been obtained from the area containing the smaller needle and the matrix (Figure 31) showed a possible preferred orientation relationship of [100]α'"
[101]γ and (011)α'" (111)γ between the martensitic lath and expanded austenite matrix in this sample, indicating closed-packed planes in each phase should be parallel to each other and it can possibly indicate the formation of coherent martensitic laths with the expanded austenite matrix and this is matched with a Nishiyama-Wasserman (N-W) orientation relationship [56].
The electron diffraction patterns obtained from two different areas in the matrix, one close to the smaller needle near the surface of sample, and the other electron diffraction pattern obtained
96! from an area close to the central (big) needle are matched with expanded γ-iron at the [110] zone axis. The calculated lattice parameter from these two diffraction patterns are (0.364 ± 0.001) nm, which corresponds to 2.3 % lattice expansion. This obtained value is matched with the calculated lattice parameter for nitrogen austenite from the XRD result of the quenched sample.
5.4.2! Cryomilled sample
Martensitic laths were the main features observed in the cryomilled sample. This is consistent with XRD results indicating that the sample is composed of mostly nitrogen martensite and around 9% of nitrogen austenite. As shown in Figure 32, the retained austenite was also transformed into lathes in this sample and came between martensitic lathes. This morphological transition of nitrogen austenite into lathes after cryomilling can possibly be ascribed as complex shear interaction during martensitic transformation as has been observed by other researchers. As can be seen in Figure 32, there is also a high density of dislocations around retained austenite lathes which can possibly arise from accommodation of the high volume misfit between transition nitrogen austenite and martensitic matrix.
The electron diffraction patterns taken from lathes in this sample possibly matched with
[113] zone axis of α'-Fe8N and [111] zone axis of α'-Fe8N. The diffuse spots in acquired electron diffraction pattern also can be due to local lattice distortion because of insertion of nitrogen atoms. However, the electron diffraction pattern obtained from small lath-shaped features also marked on Figure 35 matched with the [100] zone axis of α-iron, possibly indicating nitrogen martensitic regions with low nitrogen content. Therefore, there is no appreciable tetragonality observed in the acquired electron diffraction pattern from this area. The streaking observed in this electron diffraction pattern can possibly be accounted for grain rotation introduced by 97! interstitial nitrogen atoms or it can possibly be due to the formation of extremely small coherent precipitations in a nitrogen martensite matrix [57].
5.4.3! Heat-treated sample
From the XRD results, a composition of retained austenite, nitrogen martensite and possibly ordered nitrogen martensite (α''-Fe16N2) is expected for this sample. The STEM image shows the formation of surface martensitic needles in this sample (Figure 36) having a contrast between inner part and outer part by increasing the camera length possibly arising from two different orientations or two different phases. This contrast is also observed in the STEM image taken from the left side of the heat-treated sample by increasing the camera length. Bright field images taken from this sample also revealed the formation of three sets of martensitic needles in this sample. Each set of these needle has a different size which is parallel to each other in each set.
Different sizes of formed martensitic needles in this sample can be interpreted due to the differences in the growth rates of each set possibly depending on the orientation of a needle and on prior austenite deformation [59] [60].
The electron diffraction patterns obtained in two different zone axes of the [100] and the
[103] from formed needles in this sample match with simulated diffraction pattern of α''-Fe16N2.
Dark field images acquired from the superlattice reflections corresponding to {200} and {110} planes in the electron diffraction pattern of α''-Fe16N2 for the [100] zone axis showed that martensitic surface needles possibly are composed of two different phases. Areas near the surface of these needles light up in dark field images corresponding to these planes and these areas are consistent with outer brighter areas of needles in STEM image (Figure 36) for this sample. However, it seems these needles can possibly have different central structure (core 98! structure) as the central areas remained dark. The central parts can possibly correspond to untransformed regions of nitrogen martensite (α') to α''-Fe16N2. The solved electron diffraction pattern obtained from the matrix possibly matched with electron diffraction pattern of α'-Fe8N for the [111] zone axis. Therefore, α''-Fe16N2 has been formed in a nitrogen martensite (α') matrix in this sample. Wallace et al. and Metzger et al. both produced an α''-Fe16N2 phase as a mixture with α' and γ phases, while Coey and his co-workers produced α''-Fe16N2 as a mixture with α', γ' phases and Li et al. reported on the formation of α''-Fe16N2 in an α matrix. The formation of α''-Fe16N2 in a nitrogen martensite (α') matrix is coherent because the misfit between α''-Fe16N2 and the nitrogen martensite (α') matrix is relatively small and according to
Gibbs free energy calculations, formation of coherent α''-Fe16N2 in a nitrogen martensite (α') matrix is favorable while in a ferrite matrix, formation of a nitrogen-deficient α'' is favorable.
During tempering, the formation of stoichiometric Fe16N2 is initially expected, while structural vacancies may have introduced and formed Fe16N2-x upon prolonged tempering and form nitrogen-deficient α'' which can be due to the change of the matrix from martensite to ferrite [60,
61].
As observed in obtained dark field images from superlattice reflections, bright areas can possibly indicate formation of α''-Fe16N2 with long range ordering. Therefore, it can be concluded that these primary needles have been partially transformed into the ordered martensitic phase (α''-Fe16N2) though they did not fully transform to ordered martensitic phase of
α''-Fe16N2. This also matches what Jack reported in 2000 according to the XRD results [14].
Regarding the existence of a degree of ordering and disordering of transformation of α'-Fe8N to
α''-Fe16N2. He believes this can be a reason for discrepancies of magnetic results reported for α''- 99! Fe16N2 [14]. Also, some parts of the matrix that light up correspond to brighter regions in the
STEM image on the left side of the sample. In the dark field image which has obtained from
{200} planes, areas around the needles which have been primarily numbered as set two have lit up. The interesting point is that the central areas in all three sets have almost the same size in this dark field image. The electron diffraction pattern obtained from another part of matrix near these needles in the [111] zone axis possibly matched with α'-Fe8N. This observation can possibly confirm the Jack hypothesis about the existence of a pre-precipitation step in the transformation of disordered nitrogen martensite to ordered nitrogen martensite (α''-Fe16N2) [14]. In this step, the disordered nitrogen martensite clusters together to reach the tetragonality of α'-Fe8N and then the nitrogen atoms become ordered and form α''-Fe16N2. This is probably why the disordered nitrogen martensite close to the area containing the α''-Fe16N2 phase has the composition of α'-
Fe8N. Also, the electron diffraction pattern which has been obtained in the [103] zone axis of α''-
Fe16N2 clearly shows the superlattice reflections indicating the formation of the ordered phase.
The dark field image corresponding to {301} planes shows that two needles became bright and these two needles did not have any central dark area. Also, a group of needles labeled as series three appeared bright in this dark field image, which indicates these two sets have the same structure, thus supporting the primary hypothesis about partially transformation of disordered nitrogen martensite to ordered nitrogen martensite in formed needles. Therefore, it seems there is a degree of ordering and disordering; some of these needles have fully transformed to α''-Fe16N2, some of them got partially ordered. However, according to the proposed method in Figure 13, it is expected to form a sample containing 100% of α''-Fe16N2. 11.1 at.% of nitrogen should be primarily inserted into nitrogen austenite for forming 100% of α''-Fe16N2 after heat-treatment.
The maximum theoretical nitrogen concentration in nitrogen austenite is 10.3 at.% and the 100! maximum nitrogen content for the quenched sample is (9.8 ± 0.4) at.% , so it is not expected that all nitrogen martensite transforms into α''-Fe16N2. Novel methods such as adding 1 at.% of chromium have been proposed for trapping the nitrogen atoms, increasing the nitrogen content in nitrogen austenite beyond the maximum theoretical value of 10.3 at.% and reaching the 11.1 at.% of nitrogen content for forming 100% of α''-Fe16N2 after heat-treatment, but these results are not included in this study.
From the point of view of morphology of α''-Fe16N2, it has been reported in early microstructural studies on heat-treated steel that α''-Fe16N2 has a plate-shaped or disc-shaped morphology with a characteristic “rosette” contrast when viewed normal to the plate surface or undulated lath-shaped morphology [61]. However, TEM observations which have been done in this study on the heat-treated sample show that α''-Fe16N2 has a needle- or lath-shaped appearance and this is closer to Liu’s report that these precipitations have a ribbon-shaped morphology [22].
101! Chapter 6 -! Conclusion
Structural analysis, including TEM and XRD, has been applied to characterize iron nitrided powders produced by Jack’s route and using nitrogen austenite as a precursor for producing α''-
Fe16N2. Also, magnetic properties of these samples have been studied using VSM by S. Lan.
After the nitridation process, the sample contains nitrogen austenite with the lattice expansion of 2.3% and the nitrogen content of (9.8 ± 0.4) at.% according to XRD data and a small fraction of a magnetic phase according to VSM data. In the TEM study of this sample, a few needle-shaped coherent precipitations of nitrogen martensite with a possible preferred
orientation relationship of [100]α'" [101]γ and (011)α'" (111)γ to the expanded austenite matrix matching a Nishiyama-Wasserman (N-W) orientation relationship have been observed. These needle-shaped precipitations can be stress-induced martensite or related areas with low nitrogen content, which would imply that Ms (martensite starting temperature) had an increase; therefore, after quenching, a very small volume fraction of nitrogen martensite had been formed. A possible local rearrangement of nitrogen atoms resulting in ordering of nitrogen atoms in formed nitrogen martensites also have been observed, possibly due to the point that in these lath martensite, because of the stress fields around individual dislocation and cell walls, certain interstice positions near these defects can provide lower-energy sites for interstitial atoms such as nitrogen rather than the normal interstitial lattice positions and a local rearrangement of nitrogen atoms near to these positions occurred [59, 60].
After cryomilling the sample at liquid nitrogen temperature, the sample contains around 90% of nitrogen martensite and around 9% of retained austenite with the same lattice parameter of its parent expanded austenite, according to XRD data for this sample. Nitrogen martensite with 102! lath-shaped morphology formed in this sample according to the TEM observation. The retained austenite was also transformed into lathes and came between martensitic lathes in this sample with a high density of dislocations around retained austenite lathes which can possibly arise from accommodation of the high volume misfit between transition nitrogen austenite and martensitic matrix and comes between martensitic lathes. This morphological transition of nitrogen austenite into lathes possibly can be ascribed as complex shear interaction during martensitic transformation [60].
After heat-treating, the sample is composed of retained austenite, nitrogen martensite and possibly ordered nitrogen martensite (α''-Fe16N2). The increase in the mass magnetization at 1.5
× 107/4π (A/m) (15 kOe) from 158.4 A.m2/kg for the cryomilled sample to 202.5 A.m2/kg for the heat-treated sample can also confirm the formation of ordered nitrogen martensite (α''-Fe16N2) from disordered nitrogen martensite (α') in this sample. However, the mass magnetization at 1.5
× 107/4π (A/m) (15 kOe) of the final heat-treated sample is lower than mass magnetization at 1.5
× 107/4π (A/m) (15 kOe) for the primary α-Fe sample due to the existence of retained austenite of around 9 at.% in this sample. Coercivity also increased after heat-treating the sample in comparison to primary α-Fe powder from 5.0 × 103/4π Α/m to 110.4 × 103/4π for the cryomilled sample and to 165.1 ×103/4π Α/m for the heat-treated sample. This significant increase in coercivity is a very desirable property for a promising permanent magnetic material making the heat-treated sample containing the α''-Fe16N2 phase a candidate as it indicates that it is stable against demagnetization. In addition, the remanent magnetization increased from 0.32 A.m2/kg for the primary α-Fe sample to 5.01 A.m2/kg for the cryomilled sample after formation of the disordered nitrogen martensite and finally to 8.49 A.m2/kg for the final heat-treated sample,
103! which shows the remained magnetization after the external magnetic field is removed is higher for the heat-treated sample in comparison to the primary α-Fe sample.
From microstructural point of view, three sets of surface martensitic needles with different sizes due to the possible differences in the growth rate of each set possibly depending on the orientation of a needle and on prior austenite deformation formed after heat-treating the sample.
The electron diffraction patterns obtained in two different zone axes of the [100] and the
[103] from formed needles in this sample match with simulated diffraction pattern of α''-Fe16N2.
Dark field images acquired from the superlattice reflections in the electron diffraction pattern of
α''-Fe16N2 for the [100] zone axis showed that martensitic surface needles possibly are composed of two different phases. The central parts can possibly correspond to untransformed regions of nitrogen martensite (α') to α''-Fe16N2. Therefore, α''-Fe16N2 has been formed in a nitrogen martensite (α') matrix in this sample. The formation of α''-Fe16N2 in a nitrogen martensite (α') matrix is coherent because the misfit between α''-Fe16N2 and a nitrogen martensite (α') matrix is relatively small and according to Gibbs free energy calculations, formation of coherent α''-
Fe16N2 in a nitrogen martensite (α') matrix is favorable.
Dark field images from superlattice reflections indicate the partial transformation to ordered martensitic phase (α''-Fe16N2) with long range ordering on the surface of these needles. This observation also matches with Jack’s report regarding the existence a degree of ordering and disordering of transformation of α'-Fe8N into α''-Fe16N2, indicating not all nitrogen atoms in formed disorder nitrogen martensite get order after heat-treatment. This partial transformation and rearrangement of nitrogen atoms can possibly describe the discrepancies of magnetic results
104! reported for α''-Fe16N2. This observation can possibly confirm the Jack hypothesis about the existence of a pre-precipitation step in the transformation of disordered nitrogen martensite to ordered nitrogen martensite (α''-Fe16N2) [14]. In this step, the disordered nitrogen martensite clusters together to reach the tetragonality of α'-Fe8N and then the nitrogen atoms become ordered and form α''-Fe16N2. A possible hypothesis for this observation is that during this transformation, the areas near to the surface of the needles suck nitrogen from the central area until the average of nitrogen concentration will be equal to α'-Fe8N and then the ordering takes place. Therefore, it is expected that the central untransformed region of needles corresponding to darker areas in dark field images are areas where might be deficiency in nitrogen with non- uniform nitrogen distribution in it.
However, it is expected to form a sample containing 100% of α''-Fe16N2. 11.1 at.% of nitrogen should be primarily inserted into nitrogen austenite for forming 100% of α''-Fe16N2 after heat-treatment. The maximum theoretical nitrogen concentration in nitrogen austenite is 10.3 at.% and the maximum nitrogen content for the quenched sample is (9.8 ± 0.4) at.%, so it is not expected that all nitrogen martensite transforms into α''-Fe16N2.
105! Chapter 7 -! Future Work
If possible, an attempt should be made to prepare a thinner TEM sample with thickness of less than 30 nm, in order to apply some techniques such as quantitative electron energy loss spectroscopy (EELS) on a heat-treated sample and find the nitrogen concentration in the untransformed central areas of needles. Doing elemental mapping of nitrogen by using EELS can also reveal the relative nitrogen deficiency in untransformed regions in comparison to transformed regions of α''-Fe16N2. By preparing a TEM sample with proper thickness for high resolution transmission electron microscopy (HRTEM) with the least thickness of an amorphous layer on its surface, some observations can be done on the sharp interface between transformed regions and untransformed regions in a heat-treated sample. Differential phase contrast (DPC) imaging also can be done to visualize the variant in electromagnetic field in a heat-treated sample which can be a very useful technique for imaging the relative magnetic field around precipitations and also within a precipitation in samples, specifically the heat-treated sample. It is also suggested to do in-situ TEM and do the heat-treatment by using a heating TEM sample holder in an aberration corrected microscope, which will provide a deeper understanding of how the ordering of nitrogen atoms take place during heat-treatment.
106! References
[1] O. Gutfleisch, "Controlling the properties of high energy density permanent magnetic materials by different processing routes," Journal of Physics D: Applied Physics, vol. 33, p. R157, 2000.
[2] J. D. Livingston, "The history of permanent-magnet materials," JOM, vol. 42, pp. 30-34, 1990.
[3] A. Magnetic Materials Producers, Standard specifications for permanent magnet materials: Magnetic Materials Producers Association, 1996.
[4] S. Kozawa, "Trends and Problems in Research of Permanent Magnets for Motors- Addressing Scarcity Problem of Rare Earth Elements," Science and Technology Trends, Quarterly Review, vol. 38, pp. 40-54.
[5] T. Ogawa, "Iron Nitride Interstitial Compound: New Candidate for Permanent Magnetic Material with Rare Earth Element Less/Free."
[6] C. Hurst, "China's Rare Earth Elements Industry: What Can the West Learn?," DTIC Document.
[7] K. J. Strnat, "Modern permanent magnets for applications in electro-technology," Proceedings of the IEEE, vol. 78, pp. 923-946, 1990.
[8] H. W. F. Sung and C. Rudowicz, "Physics behind the magnetic hysteresis loops: a survey of misconceptions in magnetism literature," Journal of magnetism and magnetic materials, vol. 260, pp. 250-260, 2003.
[9] H. A. Wriedt, N. A. Gokcen, and R. H. Nafziger, "The Fe-N (iron-nitrogen) system," Bulletin of Alloy Phase Diagrams, vol. 8, pp. 355-377, 1987.
[10] D. H. Jack and K. H. Jack, "Invited review: Carbides and nitrides in steel," Materials Science and Engineering, vol. 11, pp. 1-27, 1973.
[11] K. H. Jack, "The Iron-Nitrogen System: The Preparation and the Crystal Structures of Nitrogen-Austenite and Nitrogen-Martensite," in Proceedings of the Royal Society of London A: Mathematical, Physical and Engineering Sciences, 1951, pp. 200-215.
[12] K. H. Jack, "The synthesis, structure, and characterization of α''-Fe16N2," Journal of Applied Physics, vol. 76, pp. 6620-6625, 1994.
[13] R. C. Ruhl and M. Cohen, "Splat quenching of iron-carbon alloys," Trans Met Soc AIME, vol. 245, pp. 241-251, 1969.
107! [14] K. H. Jack, "α''-Fe16N2: A Giant Magnetic Moment Material?," in Materials Science Forum, 2000, pp. 91-98.
[15] E. J. Mittemeijer, "Fundamentals of Materials Science: The Microstructure and Property Relationship Using Metals as Model Systems."
[16] L. Cheng, A. Böttger, T. H. De Keijser, and E. J. Mittemeijer, "Lattice parameters of iron-carbon and iron-nitrogen martensites and austenites," Scripta Metallurgica et Materialia, vol. 24, pp. 509-514, 1990.
[17] J. M. D. Coey, K. O'Donnell, Q. Qinian, E. Touchais, and K. H. Jack, "The magnetization of alpha," Journal of Physics: Condensed Matter, vol. 6, p. L23, 1994.
[18] K. H. Jack, "The Occurrence and the Crystal Structure of α''-Fe16N2 Iron Nitride; a New Type of Interstitial Alloy Formed during the Tempering of Nitrogen-Martensite," in Proceedings of the Royal Society of London A: Mathematical, Physical and Engineering Sciences, 1951, pp. 216-224.
[19] J. M. D. Coey and P. A. I. Smith, "Magnetic nitrides," Journal of magnetism and magnetic materials, vol. 200, pp. 405-424, 1999.
[20] H. Tanaka, S. Nagakura, Y. Nakamura, and Y. Hirotsu, "Electron crystallography study of tempered iron-nitrogen martensite and structure refinement of precipitated α''-Fe16N2," Acta materialia, vol. 45, pp. 1401-1410, 1997.
[21] J. M. D. Coey, "The magnetization of bulk α''-Fe16N2," Journal of Applied Physics, vol. 76, pp. 6632-6636, 1994.
[22] Z. Q. Liu, Y. X. Chen, D. X. Li, Z. K. Hei, and H. Hashimoto, "Microstructural investigation on the precipitation of iron nitrides and nitride transformation in ion- nitrided pure iron," Metallurgical and Materials Transactions A, vol. 32, pp. 2681-2688, 2001.
[23] T. K. Kim and M. Takahashi, "New magnetic material having ultrahigh magnetic moment," Applied Physics Letters, vol. 20, pp. 492-494, 1972.
[24] M. Takahashi and H. Shoji, "α''-Fe16N2 problems: giant magnetic moment or not," Journal of magnetism and magnetic materials, vol. 208, pp. 145-157, 2000.
[25] K. Nakajima and S. Okamoto, "Large magnetization induced in single crystalline iron films by high-dose nitrogen implantation," Applied Physics Letters, vol. 56, pp. 92-94, 1990.
108! [26] Y. Sugita, K. Mitsuoka, M. Komuro, H. Hoshiya, Y. Kozono, and M. Hanazono, "Giant magnetic moment and other magnetic properties of epitaxially grown α''-Fe16N2 single crystal films," Journal of Applied Physics, vol. 70, pp. 5977-5982, 1991.
[27] Y. Sugita, H. Takahashi, M. Komuro, K. Mitsuoka, and A. Sakuma, "Magnetic and Mössbauer studies of single crystal α''-Fe16N2 and Fe8N martensite films epitaxially grown by molecular beam epitaxy," Journal of Applied Physics, vol. 76, pp. 6637-6641, 1994.
[28] Y. Sugita, H. Takahashi, M. Komuro, M. Igarashi, R. Imura, and T. Kambe, "Magnetic and electrical properties of single phase, single crystal α''-Fe16N2 films epitaxially grown by molecular beam epitaxy," Journal of Applied Physics, vol. 79, pp. 5576-5581, 1996.
[29] M. Komuro, Y. Kozono, M. Hanazono, and Y. Sugita, "Epitaxial growth and magnetic properties of Fe16N2 films with high saturation magnetic flux density," Journal of Applied Physics, vol. 67, pp. 5126-5130, 1990.
[30] H. Takahashi, K. Mitsuoka, M. Komuro, and Y. Sugita, "Ferromagnetic resonance studies of Fe16N2 films with a giant magnetic moment," Journal of Applied Physics, vol. 73, pp. 6060-6062, 1993.
[31] C. Gao, W. D. Doyle, and M. Shamsuzzoha, "Quantitative correlation of phase structure with the magnetic moment in rf sputtered Fe-N films," Journal of Applied Physics, vol. 73, pp. 6579-6581, 1993.
[32] D. C. Sun, E. Y. Jiang, M. B. Tian, C. Lin, and X. X. Zhang, "Epitaxial single crystal Fe16N2 films grown by facing targets sputtering," Journal of Applied Physics, vol. 79, pp. 5440-5442, 1996.
[33] M. A. Brewer, K. M. Krishnan, and C. Ortiz, "Epitaxial Fe16N2 films grown on Si (001) by reactive sputtering," Journal of Applied Physics, vol. 79, pp. 5321-5323, 1996.
[34] M. Takahashi, H. Shoji, H. Takahashi, T. Wakiyama, M. Kinoshita, and W. Ohta, "Synthesis of α''-Fe16N2 films by using reactive plasma," Magnetics, IEEE Transactions on, vol. 29, pp. 3040-3045, 1993.
[35] M. Takahashi, H. Shoji, H. Takahashi, H. Nashi, T. Wakiyama, M. Doi, and M. Matsui, "Magnetic moment of α''-Fe16N2 films," Journal of Applied Physics, vol. 76, pp. 6642- 6647, 1994.
[36] M. Q. Huang, W. E. Wallace, S. Simizu, A. T. Pedziwiatr, R. T. Obermyer, and S. G. Sankar, "Synthesis and characterization of Fe16N2 in bulk form," Journal of Applied Physics, vol. 75, pp. 6574-6576, 1994.
[37] W. E. Wallace and M. Q. Huang, "Enhanced Fe moment in nitrogen martensite and Fe16N2," Journal of Applied Physics, vol. 76, pp. 6648-6652, 1994. 109! [38] X. Bao, R. M. Metzger, and W. D. Doyle, "Synthesis of high moment and high coercivity iron nitride particles," Journal of Applied Physics, vol. 73, pp. 6734-6736, 1993.
[39] R. M. Metzger, X. Bao, and M. Carbucicchio, "Magnetism of α''-Fe16N2" Journal of Applied Physics, vol. 76, pp. 6626-6631, 1994.
[40] D. C. Sun, E. Y. Jiang, and D. Sun, "Study on the structure of Fe16N2 single-crystal films by transmission electron microscopy," Thin solid films, vol. 298, pp. 116-121, 1997.
[41] A. Sakuma, "Electronic and magnetic structure of iron nitride, Fe16N2" Journal of Applied Physics, vol. 79, pp. 5570-5575, 1996.
[42] W. Y. Lai, Q. Q. Zheng, and W. Y. Hu, "The giant magnetic moment and electronic correlation effect in ferromagnetic nitride Fe16N2," Journal of Physics: Condensed Matter, vol. 6, p. L259, 1994.
[43] P. Hayes, W. Roberts, and P. Grieveson, "A hvem study of the precipitation and dissolution of iron nitrides in ferrite," Acta Metallurgica, vol. 23, pp. 849-854, 1975.
[44] X. C. Xiong, A. Redjaïmia, and M. Gouné, "In situ transmission electron microscopy investigations of the kinetics of α''-Fe16N2 precipitation during the ageing of nitrogen ferrite," Scripta Materialia, vol. 63, pp. 1232-1235.
[45] Y. Watanabe, H. Nozaki, M. Kato, and A. Sato, "Formation of Fe16N2 and development of magnetic anisotropy in a Fe-N alloy," Acta Metallurgica et Materialia, vol. 39, pp. 3161-3166, 1991.
[46] P. Ferguson and K. H. Jack, "The tempering of Fe-C¸N martensite," Scripta Metallurgica, vol. 18, pp. 1189-1194, 1984.
[47] G. Hinojosa, J. Oseguera, O. Salas, and P. S. Schabes-Retchkiman, "Microscopical characterization of α''-Fe16N2 obtained by plasma assisted processes," Scripta Materialia, vol. 34, pp. 141-145, 1996.
[48] E. Lehrer, "The equilibrium, iron-hydrogen-ammonia," Z. Electrochem, vol. 36, pp. 383- 392, 1930.
[49] J. W. Edington, "Practical Electron Microscopy in Materials Science. 2. Electron Diffraction in the Electron Microscope," Macmillan Press, x+ 122, A 4, illustrated, 1975, 1975.
[50] J. W. Edington, "Electron Diffraction in Electron Microscope," Practical Electron Microscopy in Materials Science\ Book\, vol. 2, p. 122, 1975.
[51] K. H. J. Buschow and F. R. Boer, Physics of magnetism and magnetic materials: Springer, 2003. 110! [52] B. C. Dodrill and L. S. Cryotronics, "Magnetic Media Measurements with a VSM," Lake Shore cryotronics, Inc, vol. 575, 2009.
[53] S. Nicolopoulos, D. Bultreys, and C. I. E. C. Comission, "Precession electron diffraction and TEM applications."
[54] J. Portillo, E. F. Rauch, S. Nicolopoulos, M. Gemmi, and D. Bultreys, "Precession electron diffraction assisted orientation mapping in the transmission electron microscope," in Materials science forum, pp. 1-7.
[55] J. S. Bowles and C. M. Wayman, "The bain strain, lattice correspondences, and deformations related to martensitic transformations," Metallurgical Transactions, vol. 3, pp. 1113-1121, 1972.
[56] B. V. N. Rao, "On the orientation relationships between retained austenite and “lath” martensite," Metallurgical and Materials Transactions A, vol. 10, pp. 645-648, 1979.
[57] D. Wang, "Low-temperature gas-phase carburizing and nitriding of 17-7 PH stainless steel" Case Western Reserve University.
[58] Q. Xu, V. Kumar, T. de Kruijff, J. Jansen, and H. W. Zandbergen, "Preparation of TEM samples for hard ceramic powders," Ultramicroscopy, vol. 109, pp. 8-13, 2008.
[59] E. J. Mittemeher, L. Cheng, P. J. Van der Schaaf, C. M. Brakman, and B. M. Korevaar, "Analysis of nonisothermal transformation kinetics; tempering of iron-carbon and iron- nitrogen martensites," Metallurgical Transactions A, vol. 19, pp. 925-932, 1988.
[60] M. J. Van Genderen, A. Böttger, and E. J. Mittemeijer, "Formation of α''iron nitride in FeN martensite: Nitrogen vacancies, iron-atom displacements, and misfit-strain energy," Metallurgical and Materials Transactions A, vol. 28, pp. 63-77, 1997.
[61] A. Van Gent, F. C. Van Doorn, and E. J. Mittemeijer, "Crystallography and tempering behavior of iron-nitrogen martensite," Metallurgical Transactions A, vol. 16, pp. 1371- 1384, 1985.
111!