Progress in research and development on MAX phases: a family of layered ternary compounds

Abstract

The MAX phases are a group of layered ternary compounds with the general formula M_n+1AX_n (M: early transition metal; A: group A element; X: C and/or N; n = 1–3), which combine some properties of metals, such as good electrical and thermal conductivity, machinability, low hardness, thermal shock resistance and damage tolerance, with those of ceramics, such as high elastic moduli, high temperature strength, and oxidation and corrosion resistance. The publication of papers on the MAX phases has shown an almost exponential increase in the past decade. The existence of further MAX phases has been reported or proposed. In addition to surveying this activity, the synthesis of MAX phases in the forms of bulk, films and powders is reviewed, together with their physical, mechanical and corrosion/oxidation properties. Recent research and development has revealed potential for the practical application of the MAX phases (particularly using the pressureless sintering and physical vapour deposition coating routes) as well as of MAX based composites. The challenges for the immediate future are to explore further and characterise the MAX phases reported to date and to make further progress in facilitating their industrial application.

Introduction

Metallic materials are typically characterised by being thermally and electrically conductive, plastically deformable at room temperature, readily machinable, thermal shock resistant, damage tolerant, relatively soft, etc. On the other hand, ceramics are generally characterised by high elastic moduli, good high temperature mechanical properties, good oxidation and corrosion resistance, etc. The relatively recent discovery of a family of new materials, namely the MAX phases1 (also termed in some publications ‘metallic ceramics’2) has provided materials that possess a useful combination of both metallic and ceramic characteristics.

The term ‘M_n+1AX_n phase’ was used for the first time by Barsoum in 2000.1 In the general formula M_n+1AX_n, later abbreviated to MAX,3 M represented an early transition metal, A represented an A group element, X represented C and/or N, and n = 1–3. The labelling of A group elements in the periodic table came from the old CAS notation: today, the MAX phase A-group elements are drawn from columns 13–16 in the periodic table. The transition metal, M, elements that form MAX phases are: Sc, Ti, V, Cr, Zr, Nb, Mo, Hf and Ta; the A-group elements are Al, Si, P, S, Ga, Ge, As, In, Sn, Tl and Pb. In an isolated case, Cd was also reported to form MAX phase by taking the position of A. Depending on the value of n, the M₂AX, M₃AX₂ and M₄AX₃ phases are usually referred to 211, 312 and 413 phases respectively. Experimental or calculated evidence has also been advanced for the existence of higher order MAX phases such as 514,4 615,5 and 716.6 Table 1 summarises basic data for the MAX phases that have been experimentally synthesised, alongside information on phases that have been proposed on the basis of structural observations at the unit cell level and/or theoretically predicted from calculations.

Table 1.

MAX phases reported in literature as having been experimentally synthesised, proposed and/or theoretically predicted, and some fundamental properties*: solid solutions, except those without end members reported, are not listed; polymorphs calculated as unstable are not listed

		Density	Lattice parameters		Electrical		Thermal			Mechanical
No.	Compound	d/g cm⁻³	a/Å	c	ρ ₀/μΩ m	α/K⁻¹	CTE/×10⁻⁶ K⁻¹	κ/W m⁻¹ K⁻¹	C _p/J mol⁻¹ K⁻¹	E/GPa	B/GPa	HV/GPa	σ _c/MPa	σ _f/MPa	Comments
	211 phases
1	Sc₂AlC	2·99†	3·280†	15·373†12							99†12				Calculated
2	Sc₂GaC	3·93†	3·253†	15·813†12							96†12				Calculated
3	Sc₂InC	4·72†	3·272†	16·452†12							93†12				Calculated
4	Sc₂TlC	6·60†	3·281†	16·530†12							90†12				Calculated
5	Ti₂AlC	4·111	3·051	13·63751	0·36	0·003552	8·251	4652	7854	27753	14453, 18656, 137†11, 166†55	4·551	54051
							7·1(a), 10(c)51	3353		305†55
6	Ti₂AlN	4·311	2·989	13·61451	0·2551		8·857			288†58	16956, 154†57	451	38151
							10·6(a), 9·75(c)57			28659
7	Ti₂SiC	4·35†	3·052†	12·873†11							167†11, 205†4				Predicted, stable
8	Ti₂PC	4·56†	3·191†	11·457†11							163†11				Predicted, unstable
9	Ti₂SC	4·621	3·216	11·221	0·5260		9·357	6060	8260	31661, 29060	19162, 14560	861	140061
							8·6(a), 8·7(c)57			329†60	181†63, 179†60
10	Ti₂GaC	5·531	3·07	13·521						282†64	129†11, 141†64
11	Ti₂GaN	5·751	3·00	13·31
12	Ti₂GeC	5·365	3·07	12·931	0·22765						163†11	5·565	190065
13	Ti₂AsC	5·71†	3·209†	11·925†11							175†11				Predicted, unstable
14	Ti₂CdC	9·711	3·1	14·411
15	Ti₂InC	6·366	3·134	14·07766	0·20	0·003066	9·566	26·666			121†11
					0 at 3·1 K67
16	Ti₂InN	6·541	3·07	13·971
17	Ti₂SnC	6·168	3·163	13·6791	0·22	0·003268	1068			22870	167†11	3·568
					0·07	0·00469
18	Ti₂TlC	8·631	3·15	13·981							139†11
19	Ti₂PbC	8·551	3·20	13·811							140†11
20	V₂AlC	4·071	3·1	13·831	0·2553		9·457	4853		27753	15253, 20171
							9·1(a), 10·0(c)57			308†55	175†72, 195†55
21	V₂SiC	5·20†	2·955†	11·983†10											Predicted, stable
22	V₂PC	5·381	3·077	10·911
23	V₂GaC	6·391	2·93	12·841						321†64	181†64
24	V₂GaN	5·941	3·00	13·31
25	V₂GeC	6·491	3·00	12·251			9·457				16571
							6·9(a), 15·8(c)57				191†57
26	V₂AsC	6·631	3·11	11·31			7·2(a), 14·0(c)57				204†57
27	Cr₂AlC	5·2173	2·863	12·81473	0·71	0·002873	13·373, 1253	2253	8473	27873	13853, 16571	3·573		37873
					0·7253		12·8(a), 12·1(c)57			24553	226†55, 193†72	5·274	115974	48374
										358†55				55575
28	Cr₂GaC	6·81	2·88	12·611								1·476
29	Cr₂GaN	6·82	2·875	12·771
30	Cr₂GeC	6·881	2·95	12·081			9·557			24577	18271	2·577	77077
							12·9(a), 17·6(c)57				199†57
31	Zr₂AlC	5·78†	3·2104†	14·2460†14							176†14				Calculated
32	Zr₂AlN	5·83†	3·2155†	14·2134†14							177†14				Calculated
33	Zr₂SC	6·201	3·40	12·131							18678
34	Zr₂InC	7·11	3·34	14·911
35	Zr₂InN	7·531	3·27	14·831
36	Zr₂SnC	6·968	3·3576	14·571	0·28	0·003568	8·368			17868		3·968
					0·143	0·00469
37	Zr₂TlC	9·171	3·36	14·781
38	Zri₂TlN	9·601	3·3	14·711
39	Zr₂PbC	8·268	3·38	14·661	0·36	0·014468	8·268					3·268
40	Nb₂AlC	6·501	3·10	13·81	0·29	0·002452	8·752	2352			16553, 20979	6·180		41380
							8·8(a), 6·8(c)57			28653	183†72
										299†55	205†55, 174†41
41	Nb₂PC	7·091	3·28	11·51
42	Nb₂SC_0·4	7·01†	3·27	11·41
43	Nb₂SC_x				0 at <5 K	(0·8<x⩽1·0)81
44	Nb₂GaC	7·731	3·13	13·561						323†64	196†64
45	Nb₂InC	8·31	3·17	14·371
46	Nb₂SnC	8·368	3·214	13·8021	0·45	0·002168	7·868	17·583		21668	161†57	3·868
					0·25	0·00969	6·6(a), 14·5(c)57
					0 at 7·8 K82
47	Nb₂AsC	8·0251	3·31	11·91	0 at 2 K84		7·357				224, 209†57
							2·9(a), 10·6(c)57
48	Mo₂GaC	8·791	3·01	13·181	082
49	Hf₂AlC	10·23†	3·2121†	14·3830†13							173†13				Calculated
50	Hf₂AlN	10·92†	3·1380†	14·1872†13							192†13				Calculated
51	Hf₂SC	11·36†	3·36	11·991
52	Hf₂InC	11·2466	3·309	14·72366	0·1966	0·003366	7·666	26·566			134†57
							7·2(a), 7·6(c)57
53	Hf₂SnC	11·268	3·320	14·3881	0·45	0·003468	8·168			23768		3·868
					0·42	0·002769
54	Hf₂SnN	7·721	3·31	14·31
55	Hf₂TlC	13·651	3·32	14·621
56	Hf₂PbC	11·568	3·55	14·46	0·075	0·006368	8·368					3·868
57	Ta₂AlC	11·4654	3·079	13·8605	0·26	0·004254	8·054	28·454	93·654	29254	25179	4·454	80454	36054
							9·2(a), 6·4(c)57			319†55	210†5, 221†55
											194†57
58	Ta₂GaC	13·051	3·10	13·571						367†64	217†64
	312 phases
1	Ti₃SiC₂	4·521	3·0665	17·6711	0·227	0·0032886	9·187, 9·268	3787	11087	32288	18789	490	105090	72090
					0·2185		8·6(a), 9·8(c)87	39·92		33989	187†10, 233†4
					0·232						20657, 204†57
2	Ti₃AlC₂	4·291	3·065	18·48791	0·387	0·0024586	9·091	4092		29791	16589	3·591	76493	34093
					0·35	0·003191	8·3(a), 11·1(c)57				160†72	2·593
3	Ti₃GeC₂	5·2265	3·07	17·761	0·2265		7·857	3894		34094	18694	5·065	127765
							8·1(a), 9·7(c)57				19757, 191†57
4	Ti₃SnC₂	5·99†	3·1366	18·65021							208†22
5	V₃SiC₂	5·27†	2·915†	17·535†10							219†10				Predicted, metastable
6	(V_0·5Cr.₀₅)₃AlC₂	5·31†	2·892	17·7324
7	Nb₃SiC₂	7·22†	3·13†	17·94†25							269†25				Predicted, metastable
8	Ta₃AlC₂	12·43†	3·0930	19·15923							238†5
	413 phases
1	Ti₄AlN₃	4·5895	2·988	23·37295	2·64	0·0003486	9·796	1226	15096	31026	18589	2·526	47526	35026
					2·026		9·6(a), 8·8(c)96		15826		21657, 196†57
2	Ti₄SiC₃	4·65†	3·05	22·674	0·54						245†4				Synthesised as thin film
3	Ti₄GeC₃			22·731											Synthesised as thin film
4	V₄AlC₃	5·24†	2·931	22·71945						414(x), 353(z)†97	218†97
5	Nb₄AlC₃	6·9742	3·123	24·10942	0·75	0·002542	7·242	13·542	15842	30642	214†41	2·642	51542	34642
6	α-Nb₄SiC₃ (Ref. 98)		3·1819	22·9877†						303(x), 278(z)†	241†	10·9†			Predicted, stable
7	Ta₄AlC₃	13·1835	3·092	23·70835	0·386	0·003535	8·235	38·435	18535	32435	26133	5·135	82135	37235
										362†38	266†38
8	Ti₄GaC₃	5·17†	3·0690	23·44046
	514 phase
1	Ti₅SiC₄	4·81†	3·04†	27·24†4							254†				Predicted, metastable
	615 phase
1	Ta₆AlC₅	13·69†	3·078†	34·681†5							285†				Proposed on basis of localised regions observed in TEM studies
	716 phase
1	Ti₇SnC₆	4·80†	3·2	4·16											Proposed on basis of localised regions observed in TEM studies

*ρ ₀: resistivity at 300 K; α: temperature coefficient of resistivity; CTE: coefficient of linear thermal expansion; κ: thermal conductivity; C _p: specific heat capacity at 300 K; E: Young’s modulus; B: bulk modulus; HV: Vickers hardness; σ _c: compressive strength; σ _f: flexural strength. Numerals in superscript denote references.

†Calculated value.

The MAX phases are all layered hexagonal structures with space group P6₃/mmc (no. 194). They consist of alternate near-close-packed layers of M₆X octahedra interleafed with layers of pure A atoms. The M₆X octahedra, similar to those forming in the respective MX binary carbides, are connected each other by shared edges. The main differences in the structures of the 211, 312 and 413 phases lie in the number of M layers in between every two A layers. As can be seen in Fig. 1, two, three and four M layers exist in between every two A layers in Ti₂SC, Ti₃SiC₂ and Ta₄AlC₃ crystal structures respectively. Such atomic arrangements lead to the characteristic layered structures of the MAX phases. For instance, Fig. 2a shows an SEM micrograph of a surface feature of a Ti₃SiC₂ grain growing inside a pore, where the preferential growth of the basal planes is prominent. The layered nature of the MAX phases when deformed can be better appreciated from Fig. 2b , showing deformed nanolaminates.7,8

Figure 1.

Unit cells of three representative MAX phases

Figure 2.

Crystallographic features of Ti₃SiC₂

The study of the MAX phases dates back to the work published in the early 1960s by Nowotny and his group in Vienna, who sometimes labelled them the M₂BC phases and sometimes the H-phases.9 Among the compounds reported are more than thirty 211 phases. In 2000, Barsoum listed forty-six 211 phases.1 If one ignores solid solutions, the length of this list has not changed since then, in terms of the occurrence of the 211 phases, though more have been theoretically predicted.10 ^– 14

Ti₃SiC₂, a 312 compound, is the most extensively studied MAX phase so far. In 1967, the Vienna group discovered both this compound15 and Ti₃GeC₂.16 The atypical nature of the MAX phases was first reported on by Nickl et al. in 1972,17 who found Ti₃SiC₂ to be anomalously soft for a transition metal carbide. Later, Goto and Hirai18 fabricated a 0·4 mm thick Ti₃SiC₂ polycrystalline ‘plate’ by chemical vapour deposition (CVD) and confirmed the earlier observation of Nickl et al. These early work apparently encouraged the development of bulk Ti₃SiC₂ and other MAX phases.

The number of 312 phases remained at two, until the early 1990s when Ti₃AlC₂ was discovered.19 Very recently, however, the 312 family doubled in size with the discovery of Ti₃SnC₂,20 ^– 22 Ta₃AlC₂,5,23 and (V_0·5Cr_0·5)₃AlC₂.24 The latter is quite intriguing; neither of the end members, V₃AlC₂ or Cr₃AlC₂, crystallises in the 312 configuration. Furthermore, V₃SiC₂ (Ref. 10) and Nb₃SiC₂ (Ref. 25) were predicted to be slightly metastable, by first-principles calculations. This opens the door for their possible synthesis by some non-equilibrium processes such as thin film technology.

The first 413 phase, Ti₄AlN₃, was discovered in 2000,26,27 nearly 40 years after the 211 phases, and for a while was believed to be the only one. Metastable Ti₄SiC₃ (Ref. 4 and 28–30) and Ti₄GeC₃ (Ref. 29 and 31) were found during epitaxial thin film synthesis. Bulk Ta₄AlC₃ has recently been synthesised and extensively studied.5,23,32 ^– 40 More recently, Nb₄AlC₃,40 ^– 43 V₄AlC₃,40,44,45 and Ti₄GaC₃,46 were discovered. The rapid expansion of the list of the MAX phases, particularly in the 413 phases, is encouraging and suggests that this is a fruitful line of research for the discovery of new MAX phases. It is important to note here that not all phases listed above have been fabricated in bulk, pure single-phase form.

Lin et al. 5 reported the existence of regions of Ta₆AlC₅ structure in direct atomic resolution observations of the layer stacking characteristics of TaC_x slabs and Al atomic planes in the Ta–Al–C system. Similarly, regions appearing to possess a Ti₇SnC₆ structure were recently observed in hot pressed Ti–Sn–C samples.6 Although currently relating to structures on the scale of unit cells, these results indicate the possibility that higher order MAX phases may exist.

The early review of Barsoum in 2000,1 and his introductory article3 have substantially galvanised the research field, some of the fundamental knowledge of the MAX phases can be found therein. The current status in Ti₃SiC₂ (Ref. 47) and some ternary carbides48,49 can be found in some recent reviews, for example that by Barsoum.50

In the present review, emphasis will be placed on the progress made in recent years on the synthesis and characterisation of the MAX phases.

Synthesis

Synthesis of bulk MAX phases

Early synthesis of Ti₃SiC₂ involved either reaction of gaseous reactants at temperatures as high as 2273 K,15 or by CVD.17,18 These processes, however, are not practical methods for fabricating engineering materials in the bulk form. Lis and co-workers published a series of papers over nearly 10 years, on the synthesis of bulk Ti₃SiC₂.99 ^– 103 They used combustion synthesis, and later in combination with hot isostatic pressing (HIP). However, their attempts always resulted in samples containing ancillary unwanted phases, mostly, TiC. Their best samples contained about 80–90 vol.-%Ti₃SiC₂.

While combustion synthesis is merited by its rapid reactions, in as short as a few seconds, but with sacrificed product quality, Barsoum et al. 65,104,105 employed hot pressing (HP) or HIP for the synthesis of nearly single phase bulk Ti₃SiC₂ polycrystalline samples. Such processes were soon extended to the synthesis of many other MAX phases, such as Ti₃GeC₂, Ti₂AlC, Ti₂AlN, Ti₂GeC, V₂AlC, Ta₂AlC and Nb₂AlC.65 Later, this reactive HIP process was used to fabricate predominantly single phase samples of Ti₃AlC₂.91 Zhou et al. 106 used an in situ HP solid–liquid synthesis of dense Ti₃SiC₂, that resulted in ∼92 wt-%Ti₃SiC₂.

The issue of bulk MAX phase synthesis was also tackled using pulse discharge sintering (PDS), also known as spark plasma sintering. The PDS process employed for the synthesis of the MAX phases is similar to reactive HP, however, with some advantages over the latter. The PDS process was used to fabricate bulk Ti₃SiC₂, starting with various molar ratios of reactants, such as Ti/Si/C,107 ^– 110 Ti/SiC/C,107,111 Ti/Si/TiC,112 ^– 115 Ti/SiC/TiC,116 and Ti/TiSi₂/TiC.117 In most cases, predominantly single phase, fully dense compounds were reactively sintered in as short as a few minutes. This method was also employed to synthesise Ti₃AlC₂,118 ^– 122 and Cr₂AlC,75,123 among others.

However, the reactive sintering processes outlined above, such as HP, HIP or PDS, cannot easily be applied to the mass production of MAX phase parts. Pressureless sintering from powder green compact without mechanical pressure, i.e. pressureless sintering, is much more economical and scalable.

Early attempts of pressureless sintering were first made via mechanical alloying (MA) and sintering process, which all ended up with a high content of unwanted secondary phases.124,125 Sun et al. 126 reported on the successful synthesis of high phase purity and density Ti₃SiC₂, by reactive pressureless sintering (patented). The addition of small amounts of Al was found to improve the density of the compacts and to enhance the formation of the target phase, Ti₃SiC₂.127 Fabrication of Ti₃SiC₂ by tape casting and/or cold pressing, followed by pressureless sintering of the commercial Ti₃SiC₂ powders were also reported.128,129 Such pressureless sintering processes from the respective phase pure compound powders are a promising practice for application.

Synthesis of MAX phase films

The early CVD work on the deposition of Ti₃SiC₂ films resulted in single phase samples with anomalous properties.17,18 Intriguingly, later CVD work on Ti₃SiC₂ failed to produce single phase films.130 ^– 132 This problem was mitigated in a study of the deposition of (SiC/Ti₃SiC₂)_n multilayered films.133 It follows that the use of CVD for the formation of the MAX phases has to date been limited, and only explored for Ti₃SiC₂. Problems with phase purity and process control are likely the causes for the limited interest in using CVD for MAX phases.

Much more rapid and fruitful results were obtained when the attention was turned from CVD to physical vapour deposition (PVD). Ti₃SiC₂ films were successfully synthesised by magnetron sputtering.4,28,134 ^– 136 Similar results were obtained in the Ti–Ge–C,31 Ti–Al–C,137 and Ti–Al–N (Refs. 138–141) systems. With the process of sputtering from elemental targets, MAX phase films, such as Nb₂AlC,142 V₂AlC,143 and V₂GeC,144 were also successfully fabricated, though some contained impurity phases.

In addition to these processes, pulsed cathodic arc,145 plasma sprayed coatings146 and high velocity oxy fuel spraying147 techniques were also used.

The successful thin film MAX phases synthesis has drawn much attention. However, since these films typically form preferentially at high synthesis temperatures and/or textured or single crystalline templates, their potential industrial use will be limited. On the other hand, lower growth temperatures (573 K or lower) result in nanocomposite films, that have been shown to be promising candidates for coating of electrical friction contacts owing to their metallic conductivity and good resistance to wear and corrosion.148 Rester et al. 149 successfully deposited such nanocomposite Ti–Si–C thin films by dc magnetron sputtering from a Ti₃SiC₂ target onto Si(100) and also onto a high speed steel substrate at temperatures as low as 573 K. When the deposition temperature was further decreased to 473 K, the Ti–Si–C films were deposited onto Si, polished stainless steel and high speed steel substrates, which failed to form any MAX phases.150

Following the early work on deposition of polycrystalline Cr₂AlC thin films on Si wafers,151 Walter et al. 152 proposed Cr₂AlC as the first MAX phase for which the gap from laboratory scale multiple target deposition to industrially applicable large area deposition of thin films could be bridged, though this promise has yet to be fulfilled.153 V₂GeC (Ref. 144) thin films have also, like Cr₂AlC, been sputter deposited at temperatures as low as 723 K; such low deposition temperatures are essential for commercialisation. A recent, comprehensive summary of MAX phase thin films and their properties can be found in Ref. 154.

Synthesis of MAX phase powders

The production of MAX phase powders provides an alternate route for the development of bulk single MAX phases, as well as more applications such as to MAX phase based composites.

As noted above, early work on the synthesis of Ti₃SiC₂ powders was conducted by combustion reaction process,103 with the best product containing ∼55 vol.-%Ti₃SiC₂. A solid–liquid reaction process with NaF addition155 and a fluctuation synthesis method156 were applied to synthesise Ti₃SiC₂ powders. In both cases, the Ti₃SiC₂ content in the as synthesised powders was <85 wt-%. The solid–liquid reaction process was also used to synthesise nearly single phase Ti₂SnC powders with small amounts of metallic Sn.157

Highly phase pure Ti₃SiC₂ (99·3 wt-%) powders were obtained by isothermal heating of Ti–Si–TiC powder mixtures.158 ^– 160 The nucleation and growth of Ti₃SiC₂ within TiC particles was observed (Fig. 3a ).160 The typical morphology of the particles was equiaxed with particle size of 2–4 μm (Fig. 3b ).160 The structural relationship between Ti₃SiC₂ and TiC can be understood from the high resolution Z-contrast STEM image (Fig. 3c ),161 showing the coherent interface of Ti₃SiC₂(0001)//TiC(111).

Figure 3.

Ti₃SiC₂ powder synthesis

The MAX phase powders, registered as Maxthal 312 (Ti₃SiC₂) and Maxthal 211 (Ti₂AlC), are commercially available, with the manufacturing information partly available in a patent.162

Heating of the Ti–Al–2TiC powder mixture at 1573–1473 K in flowing Ar for 15–30 min resulted in the production of 97 wt-%Ti₃AlC₂ powders, sized around 5 μm.163 However, when the elemental powders are heated, the mixture can be explosive. The latter was successfully suppressed by the addition of small amounts of Sn.164,165 Furthermore, the idea of reacting stoichiometric mixtures of Ti₂AlC and TiC to synthesise Ti₃AlC₂ powders is an interesting alternate route for the MAX phase synthesis that needs further work, especially on the economics.166

Properties

Polymorphism

The early high pressure studies were performed on some MAX phases such as Ti₃SiC₂,167,168 Ti₄AlN₃,169 Ti₂AlN and Ti₃AlC₂.168 Barsoum and co-workers168,169 were the first to report on polymorphic phase transformations in the MAX phases. They also investigated the behaviour of Ti₃SiC₂ at higher pressures.170 Further TEM171 and first-principles172 work confirmed these observations in Ti₃SiC₂. Ti₃GeC₂ was also found to show a transformation from α- to β-polymorph upon compression to 26·6 GPa, using a synchrotron radiation source.173

Polymorphism in the 211 and 312 phases can be traced to the various positions the A atoms in the unit cell can take; the M–X slabs remain unchanged.172 Polymorphism in the 413 phases is different. The more common structure is the α-structure of Ti₄AlN₃, where positions of the atoms along the c axis is ABABACBCBC (the underlined atoms represent the A-group element atoms).36 In β-Ta₄AlC₃, the layer stacking sequence is ABABABABAB.34

The issue of polymorphism in Ta₄AlC₃ is not that straight forward and sometimes controversial. In the very first report on Ta₄AlC₃, Manoun et al. 33 studied the high pressure behaviour of this compound by X-ray diffraction and found no phase transformation up to 47 GPa, assuming the α-polymorph. Discrepancies between measured and theoretical intensities were ascribed to preferred orientation. Experimental results on the structure of Ta₄AlC₃ were also reported by other groups that showed the structure to be β-Ta₄AlC₃.34,23

To account for the discrepency, Eklund et al. 36 for the first time, proposed the existence of two Ta₄AlC₃ polymorphs; Manoun et al. 33 and Etzkorn et al. 23 refined only the α-Ta₄AlC₃ structure, while Lin et al. 34 characterised β-Ta₄AlC₃. Eklund et al. 36 also supposed the structure reported by Manoun et al. 33 to be β-Ta₄AlC₃, with the evidence of close lattice parameters to the work of Lin et al. 34 However, thermodynamically, the existence of both α- and β-Ta₄AlC₃ at ambient conditions does not seem convincing.

To shed more light on this issue, Du et al. 39 studied the two polymorphs using first-principles calculations and showed that the total energy of α-Ta₄AlC₃ is ∼0·0866 Ry lower than that of β-Ta₄AlC₃, indicating that the α-polymorph is more stable (Fig. 4). The strong hybridisation between Ta1 5d and C1 2p, Ta2 5d and Al 3p in α-Ta₄AlC₃ result in the strong interaction between Ta1-C1 and Ta2-Al atoms, which stabilises the α-polymorph. Similar results were reported from different groups almost simultaneously.40,174,175

Figure 4.

Total energy versus volume of α- and β-Ta₄AlC₃ (Ref. 39)

Furthermore, a thermodynamic study of Ta₄AlC₃,40 using first principles calculations confirmed that at temperature below 1875 K the α-polymorph is more stable, and above this critical temperature a transformation to β-Ta₄AlC₃ may occur. The other 413 phases studied, Ti₄AlN₃,40 V₄AlC₃,40 and Nb₄AlC₃,40,41 all show stable α-polymorphs at temperatures up to 3000 K.

Phase stability

Ti₃SiC₂ was reported to become unstable at different temperatures ranging from 1273 to 1673 K.176 ^– 181 However, Barsoum et al. 182 showed that Ti₃SiC₂ was thermodynamically stable up to at least 1873 K in vacuum for 24 h. They further argued that the reduced temperature at which Ti₃SiC₂ decomposed as observed by others was due to the presence of impurity phases in the starting powders.183

Ti₃AlC₂ was reported by Wang et al. 184 to be stable to at least 1744 K when heated in Ar with a low oxygen content. Selective oxidation of Al in Ti₃AlC₂ took place on top of the specimen, which resulted in the transformation of Ti₃AlC₂ into substoichiometric TiC_x and Al₂O₃.

Good thermal stability was also reported in some bulk 211 phases, such as Ti₂AlN,26 Ti₂SC,61 and Ta₂AlC,185 which were found to be stable at up to 1723, 1873 and 1873 K respectively.

In general, the thermal stability of the MAX phases depends on their constitutive elements, the atmosphere and the vapour pressure of the elements, particularly the A elements, such as Si and Al. For example, it has been shown that thin films decompose at lower temperature in comparison to their bulk counterparts, such as in Ti₃SiC₂,180,181,186 and Ti₂AlN.187 This was attributed to the difference in diffusion length involved and possibly the measurement sensitivity used in the respective studies.180,181

It is worthwhile to mention the ‘Ti₃SiC₂ pest’ phenomenon.188 Ti₃SiC₂ was observed to undergo rapid and catastrophic disintegration above 1573 K in nitrogen, which was attributed to the volume change associated with the formation of mixtures of TiC_x, Ti(C,N)_x and TiN.188 A prevention process for such ‘Ti₃SiC₂ pest’ was suggested by increasing the oxygen partial pressure in the nitrogen atmosphere.

Electrical, thermal and elastic properties

The MAX phases are good electrical conductors, with reported room temperature resistivities ranging from 0·07 (Refs. 68 and 69) to 2 μΩ m.26 The low resistivities measured, for example, in Ti₂SnC have been difficult to reproduce,68 which has been attributed to compositional variation and the presence of impurity phases in the microstructure;1 this value must therefore be treated with caution. As shown in Table 1, most of the MAX phases have lower resistivities than Ti metal (∼0·4 μΩ m). To date, the electrical conductivities of only around one-third of the reported MAX phases have been measured.

Yoo et al. 189 were the first to report on the negligible Seebeck coefficient or thermopower in Ti₃SiC₂ over the wide temperature range of 300–850 K. Later that range was extended to 5 K.94 The same is true of Ti₃GeC₂ and Ti₃(Si,Ge)C₂ solid solutions.94 This unusual phenomenon was found to be related to the anisotropy of the crystal structure.190 The components of the thermoelectric tensor of Ti₃SiC₂ and Ti₃GeC₂, but interestingly enough not Ti₃AlC₂, in the basal plane and along the c axis were found to have opposite signs.191 The near-zero thermopower of Ti₃SiC₂, Ti₃GeC₂ and most likely other MAX phases, suggests that these compounds could be useful as electrodes during thermopower measurements of other compounds.

The coefficients of thermal expansion of the MAX phases range from 7·2×10⁻⁶ to 13·3×10⁻⁶ K⁻¹ (Table 1). Currently, the highest coefficients of thermal expansion were reported in Cr₂AlC (Ref. 73) and Cr₂GeC,57 which is one reason the former is a good candidate for large area coatings on steels.152 Measurements of the thermal expansion of the MAX phases in different orientations were performed using high temperature X-ray diffraction61 or synchrotron X-ray diffraction. For the most part, the MAX phases show little anisotropy in thermal expansion,61,192,193 usually with higher thermal expansions in the c than in a direction. The lower thermal expansion in a direction is apparently due to the stronger M–C bonding than the M–A bonding in the c direction.

The room temperature thermal conductivity data available to date for the MAX phases range from 12 to 60 W K⁻¹ m⁻¹. It is not surprising to find that the 413 compound Ti₄AlN₃, that shows the lowest electrical conductivity, shows the lowest thermal conductivity, due to the low carrier mobility which might be caused by the point defects in the crystal structure.194 The room temperature thermal conductivity of 60 W K⁻¹ m⁻¹,60 measured in Ti₂SC is the highest to date, despite the fact that its electrical conductivity (1·92×10⁶ Ω⁻¹ m⁻¹) is relatively poor.60 This is due to the large contribution from phonon channel in comparison to that of other MAX phases.60 A large phonon contribution to the thermal conductivity suggests a stiff material. This is consistent with the electronic structural study of Ti₂SC.63 The stiffness, as represented by the high bulk modulus, was found to originate from the strong Ti 3d–S 2p hybridisation; the unusual strong M–A bonding in the MAX phases can also be seen from charge density maps as shown in Fig. 5.63

Figure 5.

Calculated valence charge densities (eÅ⁻³) for (1120) plane of Ti₂AlC and Ti₂SC, revealing strong Ti–S interaction in Ti₂SC (Ref. 63)

The MAX phases are characterised by high elastic constants, as represented by the Young’s modulus E and bulk modulus B. Values of E for MAX phases measured to date range from 178 GPa (Zr₂SnC),68 to 340 GPa (Ti₃GeC₂),94 while B was found to range from 138 GPa (Cr₂AlC),53 to 261 GPa (Ta₄AlC₃).33 The value of B for Ta₄AlC₃ was subsequently obtained as 198 GPa from the measured Young’s and shear moduli.35 This discrepancy could be due to the different experimental methods used in the two studies; a degree of uncertainty also exists over the polymorphic structure of the specimen used in the earlier study.

On the other hand, the highest stiffness value of the 211 phases measured to date is for Ti₂SC (E = 316 GPa),61 almost equivalent to those of the stiffest 413 phases. This is, however, not due exclusively to the strong M–X bonds or a larger concentration of M–X bonds as in the 413 phases. Ti₂SC has the smallest c/a ratio and consequently, exceptionally strong M–A bonds (Fig. 5) which contributes to the high stiffness.63 This bonding also accounts for its high thermal conductivity.60 However, several 211 phases, e.g. V₂AlC,71 Nb₂AlC,79 Nb₂AsC,57 and Ta₂AlC,79 show higher B values than that of Ti₂SC. Experimental data for B show high scatter, particularly among data sets measured with different methods. Even higher discrepancies are common among B values calculated from first-principles.

These comments notwithstanding, the agreement between theory and experiment is quite good. One exception is Cr₂AlC for which the discrepancy in the theoretical and experimental results, particularly in bulk modulus which is higher than usual. Schneider et al. 195 studied the elastic properties of Cr₂AlC thin films by nanoindentation and ab initio molecular dynamics and reported the agreement between the experimental and theoretical results. The author and co-workers have found, however, that the interactions of the d-electrons in the density function theory, which have been so far ignored, play an important role in the elastic properties. The negligence of the Cr 3d on-site Coulomb energy, in the theoretical calculations, resulted in the overestimation of the bulk modulus of Cr₂AlC.196 The revised values are more in line with experiment.71

Hardness, machinability, deformation and fracture

The reported data of the hardness of the MAX phases range from 1·4 to 8 GPa (Table 1). The low hardness, together with their layered nature, endows the MAX phases with probably their signature property, namely, excellent machinability. Figure 6 shows Ti₃SiC₂ samples fabricated using PDS and machined to various shapes either by lathe cutting, surface miller, electrical discharge machining and/or precision machining.2,8

Figure 6.

Ti₃SiC₂ fabricated with reactive PDS process and machined to different geometries8

The machinability of Ti₂SC was successfully exploited in the development of Pb free free-cutting steels, as can be found in a publication in Japanese197 and some patents. In the metallurgical literature, Ti₂SC is usually referred to as Ti₄C₂S₂, since it has long been found as an inclusion phase in steels and superalloys.198 ^– 200 The strain induced precipitation of Ti₂SC and its possible self-lubricating properties199,200 are fruitful areas of the study of other MAX phases. It is interesting to note that Ti₂SC is the hardest among the MAX phases reported to date and also possesses the highest stiffness (Table 1). With the success of Ti₂SC as a free-cutting element in steels and the fact that other MAX phases are softer, one is greatly encouraged to develop more such free-cutting MAX phases to replace Pb in free-cutting stainless steels and even in most brasses, to eliminate the harmful effects of Pb on the environment.201,202

Basal plane slip in the crystal structures of the MAX phases enables dislocation activity, and hence local plastic deformation, even at room temperature. Figure 7 shows the slip lines observed on the surface of the Ti₃SiC₂ sample surface near a Vickers hardness indentation made at room temperature.2 These slip lines resemble those observed on the surface of deformed metal samples. It is this microplasticity, together with its good thermal conductivity, that is mostly responsible for the good thermal shock resistance of the MAX phases. For example, quenching polycrystalline Ti₃SiC₂ samples from 1673 K into water did not affect their post-quench flexural strengths.105 The microplasticity is also responsible for the remarkable damage tolerance of the MAX phases, such as reported for Ti₃SiC₂,203 Ti₃AlC₂,91 Ti₂SnC,204 Ta₂AlC,54 Cr₂AlC,74 Ta₄AlC₃,35 Nb₂AlC,80 Ti₄AlC₃,26 and others.

Figure 7.

Room temperature plastic deformation in Ti₃SiC₂ grain2

Like conventional ceramics, the MAX phases have higher compressive,205 than tensile206,207 or flexural strengths.208 Figure 8 shows the deformation and strength of Ti₃SiC₂ tested under compression and four-point bending conditions at various temperatures and the samples after deformation.2 Compressive strengths as high as 1·9 GPa have been reported for some MAX phases.65

Figure 8.

Mechanical properties of Ti₃SiC₂ prepared by reactive PDS process

The concept of incipient kink bands was proposed by Barsoum et al. 209,210 to understand the reversible non-linear elastic deformation in polycrystalline Ti₃SiC₂ subjected to cyclic loading. When Ti₃SiC₂, and other MAX phases were studied in the porous form, higher absolute energies were dissipated.7,211 Since kinking is a form of plastic buckling, the lack of constraint, due to the presence of the pores must facilitate the formation of more incipient and regular kink bands, which in turn lead to higher energies dissipated per volume.7 Some results are shown in Fig. 9.7,209,210 Along this line, an increase in porosity reduces the threshold stresses needed for incipient kink band formation.7,212

Figure 9.

a summary of energy dissipation data, in the form of a logarithm correlation between the irreversible work and the applied maximum stress, of recently investigated kinking nonlinear elastic solids,209,210 including the results for a 43% porous Ti₃SiC₂,7 and b the underlining mechanism revealed by kink band formation and concomitant delamination of a solid bridge between pores in Ti₃SiC₂ (Ref. 7)

The fracture toughness K _IC values of the MAX phases range from 4·5 to 16 MPa m^1/2 for Ti₃SiC₂,106,213 ^– 219 and from 4·6 to 9·1 MPa m^1/2 for Ti₃AlC₂.93,118,220 The K _IC values mainly depend on microstructure; generally coarse grained microstructures result in higher K _IC values due to the deflection and delamination of cracks by single grains. These mechanisms result in a profusion of crack bridging processes in the crack wake (Fig. 10a ).218 This ligament is essentially one that is held together by strong Ti–C bonds. This micrograph is very reminiscent of typical ligaments seen in wood.

Figure 10.

a lamellae bridge observed in crack wake of coarse grained Ti₃SiC₂ in R curve measurement218 and b 8 inch Ti₂AlC block indented by steel hammer (courtesy of 3-ONE-2, LLC)

The fracture toughness values of some other MAX phases were also reported in similar range, e.g. Nb₄AlC₃ (7·1 MPa m^1/2),42 Ta₄AlC₃ (7·7 MPa m^1/2),35 Ta₂AlC (7·7 MPa m^1/2),54 Ti₂AlC (6·5 MPa m^1/2),221 etc.

The good fracture toughness and the microplasticity results in behaviour that is very reminiscent to that of metals, as evidenced in Fig. 10b , where the surface of an 8×4 inch Ti₂AlC block was repeatedly hit with a heavy steel hammer. The latter only caused the formation of surface dents.

Fatigue and creep

The fatigue behaviour of MAX phases have been reported only in Ti₃SiC₂,216,218,222 which exhibits cyclic fatigue crack propagation behaviour similar to typical structural ceramics, i.e. a much larger dependence of crack growth rates on K _max, as compared to ΔK (Fig. 11).222 Nevertheless, Ti₃SiC₂ is characterised by higher threshold stress intensity factors and lower crack growth rates, that are superior to conventional structural ceramics.222 The high fatigue resistance originates from the heterogeneous and laminated structure. The fatigue threshold ΔK _th of Ti₃SiC₂ was found to remain essentially unchanged between 298 and 1373 K;216 however, at 1473 K there was a sharp decrease in ΔK _th, due to the formation of a large number of deformation microcracks. The crack tip shielding was found to result from both the bridging of entire grains and from deformation kinking and bridging of microlamellae within grains.

Figure 11.

Relationship between crack growth rates da/dN and a applied stress intensity range ΔK and b maximum applied stress intensity K _max, for Ti₃SiC₂ at load ratios R of 0·1 and 0·5 (Ref. 222)

Similarly, the creep behaviour of MAX phases has been also mainly reported in Ti₃SiC₂.223 ^– 225 Tensile creep tests performed on fine and coarse grained Ti₃SiC₂ microstructures in the 1273–1473 K temperature range showed primary, secondary and tertiary creep regimes.223,224 In compression, Zhen et al. 225 reported that the creep behaviour of Ti₃SiC₂ is also characterised by three regimes. The stress exponents of ∼2, in the quasi-steady state regime were comparable to those measured in tension. The creep rates in compression, however, were roughly an order of magnitude lower than in tension. With an increase either in stress and/or in temperature, a change of mechanism from dislocation creep to possibly subcritical crack growth was proposed.

With the potential use of the MAX phases as structural materials for high temperature applications, the investigation into their creep behaviour is urgently needed.

Oxidation

The oxidation of Ti₃SiC₂ in the form of powder176 or bulk226 ^– 229 has been studied and the oxidation kinetics was reported to be parabolic during exposure at 1173–1673 K.226 The oxidation rate was found to be lower than that of TiC, with oxides of TiO₂ and SiO₂ formed.176 The oxidation occurs by the outward diffusion of Ti and C and the inward diffusion of oxygen.226

A thin SiO₂ rich scale was found to form on the surface of Ti₃SiC₂ after short term oxidation at 1273 K in air.227 It is important to note, however, that long term oxidation, of Ti₃SiC₂ has shown that after ∼100 h, the oxidation kinetics revert from parabolic to linear and that the highest temperature continuous use temperature of Ti₃SiC₂ is probably 1173 K. The reason for the change in mechanism is not fully understood, but could be due to the formation of microcracks in the rutile outermost layer.230

Anomalous oxidation with higher kinetics at lower temperatures was observed in Ti₃SiC₂ when tested at 773–1073 K in flowing dry air.231 This phenomenon was believed to result from the formation of microcracks in the oxide scales due to phase changes in the oxide products. If this phenomenon is ubiquitous then it should be carefully examined. However, the authors have argued that will not be a serious problem in practice since any cracks formed were partially filled with amorphous SiO₂. This is especially true in Ti₃AlC₂ given the crack healing abilities of Ti₃AlC₂ at least during the early stages of oxidation.232

In general, Ti₃AlC₂ has much better oxidation resistance than Ti₃SiC₂. A comparison of the oxidation kinetics shows that the parabolic rate constants of Ti₃SiC₂, at the same temperatures, were 2–4 orders of magnitude greater than those of Ti₃AlC₂.93,233 The superior oxidation resistance is attributed to the formation of an adherent, continuous inner Al₂O₃ layer, which forms below an outer TiO₂ layer.233 The adherent, high integrity Ti₃AlC₂/Al₂O₃ interface was confirmed by high resolution TEM study (Fig. 12).234

Figure 12.

High resolution TEM image of Ti₃AlC₂/Al₂O₃ interface formed after oxidation at 1473 K for 8 h (Ref. 234)

Similarly, Ti₂AlC was found to be exceptionally oxidation resistant, because it forms a stable and adherent protective Al₂O₃ scale.235 The performance up to 1848 K, was evaluated in vacuum levels down to 10⁻² Pa. Ti₂AlC showed a parabolic and stable oxide growth rate up to 1673 K in air. Cyclic testing confirmed that the oxide layer is adherent to the bulk.235

The oxidation mechanism in Cr₂AlC, when investigated in 1373–1523 K temperature range, was also found to be by the outward diffusion of Al, together with small amounts of Cr, and the inward diffusion of O to form a surface layer of α-Al₂O₃. Here, carbides (Cr₇C₃ and Cr₃C₂), rather than oxides, formed in a layer under the outermost oxide layer. The oxidation of Cr₂AlC also obeyed a parabolic rate law,236,237 with the parabolic rate constants tabulated in Table 2. Comparison with two representative 312 phases indicate that Cr₂AlC shows comparable or superior oxidation resistance with Ti₃AlC₂ and much better oxidation resistance than Ti₃SiC₂.

Table 2.

Parabolic oxidation rate constants of some typical MAX phases

	Parabolic rate constants/kg² m⁻⁴ s⁻¹
Temperature	Ti₃SiC₂ (Ref. 226)	Ti₃Si_0·9Al_0·1C₂ (Ref. 238)	Ti₃AlC₂ (Ref. 233)	Cr₂AlC
1273 K	1·3×10⁻⁹; 1·0×10⁻⁹	3·2×10⁻¹¹	4·1×10⁻¹¹
1373 K	2·1×10⁻⁷ (Ref. 239)	1·2×10⁻¹⁰	2·7×10⁻¹⁰	1·1×10⁻¹² (Ref. 236)
1473 K	6·58×10⁻⁷; 1·92×10⁻⁶	4·0×10⁻¹⁰	4·2×10⁻¹⁰	6·8×10⁻¹⁰ (Ref. 237)
1523 K				7·1×10⁻¹⁰ (Ref. 236)
1573 K	2·7×10⁻⁶; 1·3×10⁻⁵	2·4×10⁻⁹	1·1×10⁻⁹
1673 K	1·4×10⁻⁵; 1·0×10⁻⁴

Corrosion

The corrosion resistance of Ti₃SiC₂ in acid or alkaline solutions240,241 has been found in some instances to be better than that of pure Ti.241 The enhanced resistance was attributed to the formation of a thin, passivating SiO₂ based layer. Passivating layers were also found to be responsible for the good corrosion resistance of Ti₃GeC₂ and Ti₂AlN.242 The electrochemical corrosion behaviour of more MAX phases has been studied in NaOH, HCl and H₂SO₄ solutions.243 Active dissolution or passivation was observed, depending on the solutions and potentials applied, but a generalised corrosion resistance diagram for the MAX phases is still lacking.

Ti₃SiC₂ was found to suffer from serious hot corrosion at temperatures above ∼1123 K when coated with or immersed in Na₂SO₄.244 ^– 246 The mass gains during the hot corrosion were greater than those of specimens oxidised in air at the same temperatures.244

By contrast, Cr₂AlC was found to exhibit exceptionally good hot corrosion resistance against molten Na₂SO₄.247,248 The parabolic rate constant for the hot corrosion of Na₂SO₄ coated Cr₂AlC at 1273 K was 3·4×10⁻¹¹ kg² m⁻⁴ s⁻¹.247 The corrosion resistance was attributed to the formation of a protective Al₂O₃ rich scale.

In line with its excellent oxidation resistance, Ti₂AlC was found to exhibit good corrosion resistance in Na₂SO₄.249 A protective continuous Al₂O₃ layer was formed that imparted good corrosion resistance; consequently, the corrosion kinetics was generally parabolic at 1123 K. However, the porous oxide scale failed to protect the Ti₂AlC substrate at 1173 and 1273 K.

In view of potential applications in contact with molten metals, Ti₂AlC and Ti₃SiC₂ were tested for corrosion in circulating molten Pb at 923 and 1073 K. The extent of reaction was minimal for both materials.250 However, molten Al was found to react with Ti₃SiC₂ to form interconnected networks of TiC_0·67 and molten Al.251

Friction and wear

Tribological studies have to date concentrated on Ti₃SiC₂, which was initially expected to be a material with ultra low friction, based on an early experimental report that the friction coefficient μ, measured on cleaved basal planes of individual grains, was 0·0025.252 When sputtered Ti₃SiC₂(0001) thin films were measured at normal loads of 0·1 mN to 0·24 N; however, μ of 0·1 and 0·8 were obtained for low and high applied loads respectively.253 The high friction and fast wear were attributed to excessive debris resulting in third-body abrasion.

Two stage friction and wear behaviour was reported in Ti₃SiC₂,254 ^– 256 and Ti₃AlC₂, as well as its composites,257 when tested against various counterparts. For example, El-Raghy et al. 254 found that the friction of Ti₃SiC₂, either coarse or fine grained, against a steel always underwent an initial transition period with μ of 0·15–0·45 followed by a static period with a μ of ∼0·8. The transition from the first stage to the second stage, with higher μ values and wear, was attributed to debris accumulation.254,257

Self-lubricating properties of Ti₃SiC₂ were reported against an oscillating diamond pin (μ = 0·05–0·1), but not against Ti₃SiC₂ itself.258 The low μ values of Ti₃SiC₂ against diamond were attributed to the formation of a thin film on the Ti₃SiC₂ tribosurface.

The wear rates reported for the MAX phases are usually in the order of 10⁻⁵ mm³ N⁻¹ m⁻¹ or higher,254,259 ^– 261 which is considered to be higher than the wear rate tolerable in service, i.e. 10⁻⁶ mm³ N⁻¹ m⁻¹. Exceptionally low wear rates were reported for Ti₃AlC₂ (∼2×10⁻⁶ mm³ N⁻¹ m⁻¹),262 Ta₂AlC and Ti₂AlC (⩽1×10⁻⁶ mm³ N⁻¹ m⁻¹).261 Attempts have been made to develop wear resistant composites by incorporating hard particles such as SiC,263 or Al₂O₃,257,264 into the MAX matrix. This approach was found to be effective in reducing the wear rate.

Other surface phenomena

Spontaneous growth of whiskers from a MAX phase was first reported in 1999, where Ga whiskers were observed to ‘self-extrude’ from Cr₂GaN bulk samples. Initially, it was postulated that the basal planes of Cr₂GaN were the Ga source.265 Shortly thereafter, it was postulated that the actual Ga source was excess, or unreacted, Ga present at the grain boundaries.266 Similar phenomena have also been reported in Zr₂InC,267 and in a Ti–Sn–C system.268 The driving force for the room temperature spontaneous growth of low melting point metal whiskers was attributed to the diffusion of oxygen into the grain boundaries of the substrate from which the whiskers were sprouting.267

However, the room temperature spontaneous Ga whisker growth from Cr₂GaC,269,270 could not be explained by any of the proposed mechanisms. Instead of being extruded from the grain boundaries of the Cr₂GaC, Ga whiskers were observed to grow from freshly cleaved basal planes from near surface grains. More convincing observations were made on the fracture surface of a Cr₂GaC sample, where a huge (in comparison to the grain it grew from) Ga whisker grew from the basal plane of a fractured grain (Fig. 13).269 Given that the basal plane on the fracture surface is unconstrained, mechanisms employing stress and extrusion cannot be responsible. Instead, mechanisms such as solidification (the melting point of Ga is near room temperature), and/or recrystallisation may have to be considered to explain this phenomenon. Interestingly, room temperature spontaneous growth of Ga freestanding nanoribbons from Cr₂GaC surfaces was also observed.270,271 Future work on these aspects are required to further understanding of such phenomena.

Figure 13.

Ga whisker observed on basal plane of a Cr₂GaC grain on fracture surface: whisker growth axis measured to be Ga[010] (Ref. 269)

After creep failure at 1573 K, silica based nanofibres with diameters of ∼250 nm and lengths of up to a few tens of micrometres were observed on the fracture surfaces of Ti₃SiC₂.272 The formation of a silicon oxide layer immediately after creep failure and their subsequent elongation and tearing in the crack wake was postulated to be the mechanism for the formation of these fibres. Later, synthesis of AlN nanowires by the nitridation of Ti₃Si_0·9Al_0·1C₂ solid solution was reported.273 Single crystalline AlN nanowires with the hexagonal wurtzite structure were easily prepared. These reports imply that the MAX phases are promising raw reactants for the synthesis of various one- and two-dimensional nanostructures of nitrides and oxides.

Materials and processes based on MAX phases

Surface technology

Surface modification

Owing to its several attractive properties, Ti₃SiC₂ has been considered for high temperature applications. However, compared with traditional binary carbides, its hardness and wear resistance are low. A useful approach in order to enhance the latter while maintaining the attractive properties is surface modification or the application of hard coatings.

To enhance the surface hardness, El-Raghy and Barsoum274 carburised Ti₃SiC₂ surfaces and increased the hardness from 4 GPa to 20–25 GPa. Meanwhile, silicidation of Ti₃SiC₂ enhanced the high temperature oxidation resistance by about three orders of magnitude.274 Surface boronising for Ti₃SiC₂ increased the hardness from 3·7 GPa to a maximum of 9·3 GPa.275

During the work of surface carburisation, sicilidation, composite and joining processes, Ti₃SiC₂ was found to transform to TiC by the de-intercalation of Si from the structure when the Ti₃SiC₂ is in contact with Cu,276 Ni,277 graphite,274 molten cryolite186 and liquid Al.251 Stimulated by these findings, Guo et al. 278 deposited a Cu film on Ti₃SiC₂ by magnetron sputtering followed by annealing in the 1173–1373 K temperature range. Owing to the formation of TiC, the surface hardness increased from ∼5 GPa to a maximum of ∼10 GPa, depending on the relative amount of TiC, which was controlled by the Cu film thickness.

Wan et al. 279 fabricated a hard–soft–hard sandwich structure of Al₂O₃–Ti₃SiC₂–Al₂O₃, and found the maximum strength to be 14·5% higher than that of the monolithic Ti₃SiC₂. The strong Al₂O₃/Ti₃SiC₂ interface was attributed to the formation of Si and Ti–Si liquid phases and their penetration into the Al₂O₃ layers.280

Diamond on Ti₃SiC₂

Diamond nucleation on non-diamond surfaces without pretreatment is reported to be difficult and slow. By enhancing the nucleation kinetics, such as scratching, nucleation densities of 10⁸ cm⁻² can be obtained on Si and carbide forming refractory metals, such as Nb, Mo and Ti. It has been found that a carbide interlayer is initially formed on these substrates, before diamond nucleation, such as SiC on Si, Mo₂C on Mo and TiC on Ti (e.g. Refs. 281 and 282).

On the surface of a Ti₃SiC₂ ample fabricated via the reactive PDS process, Yang et al. 282 investigated the diamond nucleation and growth by microwave plasma enhanced chemical vapour deposition. A much higher diamond nucleation density, and consequently, much higher film growth rates were obtained on Ti₃SiC₂ as compared to Si wafers. Furthermore, the diamond films on Ti₃SiC₂ were smoother with a preferred [0001] orientation and exhibited better adhesion than those on Si.282

In general, nanocrystalline diamond (NCD) films have several advantages over microcrystalline diamond (MCD) films. The high diamond nucleation density on Ti₃SiC₂ is thus a great advantage for the synthesis of NCD films. NCD films were synthesised on Ti₃SiC₂ by MPCVD under typical conditions for MCD synthesis.283 A diamond nucleation density of 10¹¹ cm⁻² was achieved and consequently dense NCD thin films with an average grain size of ∼30 nm (Fig. 14) and film growth rates of ∼2 μm h⁻¹ were obtained. Cross-sectional SEM images showed NCD films with an equiaxed grain structure and a dense film/substrate interface.

Figure 14.

Atomic force microcopy image of surface of Ti₃SiC₂ after 20 min diamond deposition283

Diamond thin films of ∼10 μm thick were also successfully deposited onto Ti₃SiC₂ with different geometries.284 The diamond coated surfaces of all specimens were polished up to a surface roughness of ∼0·2 μm. Friction coefficients between the polished diamond plate and a stainless steel rail, the plate specimen, and between a rotating rod and a stationary outer ring, were measured, and the lowest friction coefficients in both specimens were <0·01.284

It thus appears that Ti₃SiC₂ is an excellent heterosubstrate for diamond nucleation and growth. Furthermore, it can be readily machined to complicated components before the deposition. Once an adhesive diamond thin film is successfully deposited on it, the applications of diamond thin films can be greatly expanded. In the meantime, deposition of a diamond film on the surface of Ti₃SiC₂ would be of great importance to improve its tribological properties.

Joining

To further extend the application of the MAX phases to more complicated and larger dimensions, joining of the MAX phases, mainly of Ti₃SiC₂, has been investigated.

Gao et al. 285 carried out diffusion bonding for joining a Ti₃SiC₂ with a Ti–6Al–4V alloy and found the reaction was rate controlled by the solid state diffusion below 1623 K. The strength of such processed joint, however, is quite low, with a value of 100 MPa, which was attributed to the tensile stresses developed in the Ti₅SiC_x single layer.

Hashimoto et al. 286 employed slurry spin coating and reactive sintering techniques for joining Ti₃SiC₂. Slurries of Ti, Si and TiC powders were spin coated on Ti₃SiC₂ test pieces under various conditions. The coated pieces were joined together and sintered, under pressure, at temperatures of 1523–1573 K. The bonding layers were 20 or 30 μm thick and showed finer grain size than those farther from the interface.

Yin et al. 277 studied the diffusion bonding of Ti₃SiC₂ and Ni and found the growth of the reaction layer to follow a parabolic law. The diffusion of nickel through the reaction zone toward Ti₃SiC₂ was found to be the main controlling step in the bonding process. Joint strengths were measured with shear tests and showed maximum shear strengths of the order of 121 MPa, a value believed to be close to that of bulk Ti₃SiC₂.

MAX phases with similar constituent elements, such as Ti₃SiC₂–Ti₃GeC₂ and Ti₂AlC–Nb₂AlC, were found to form solid solution at interfaces by interdiffusion.287 This interdiffusion also enabled the bonding of Ti₃SiC₂ via Al interlayers,288 or bonding Ti₃AlC₂ via Si interlayers.289 In both cases, diffusion of the interlayer element into the MAX phases led to the formation of Ti₃(Si,Al)C₂ solid solutions rather than intermetallic compounds. Flexural strength reached about 65 and 80% for the Ti₃SiC₂ joint with Al interlayers and the Ti₃AlC₂ with Si interlayers of their respective monolithic MAX phases. Direct bonding of Ti₃SiC₂ and Ti₃AlC₂ by interdiffusion is also possible. The resulting bonding strengths are slightly lower than those made using interlayers.290

It is worth noting that with good machinability, parts made of MAX phases can be easily joined by mechanical methods such as threading, flanging, etc. in addition to the bonding processes discussed above.

Composites

MAX–MX and MAX–AX

Since SiC and TiC are thermodynamically stable with Ti₃SiC₂, and SiC has excellent oxidation resistance, Ti₃SiC₂–SiC and Ti₃SiC₂–TiC were the early MAX based composites synthesised and investigated.291 ^– 293 Tong et al. 291 found higher flexural strengths in Ti₃SiC₂–20SiC–8TiC (vol.-%) composite than monolithic Ti₃SiC₂, particularly at high temperatures. The fracture toughness, however, was reduced. This was attributed to a lack of large grains in the composite microstructure. The hardness was measured to be 9 GPa, a value that is close to that reported by Radhakrishnan et al. 292 for a Ti₃SiC₂–14 vol.-%SiC composite.

In the study of two composites, Ti₃SiC₂–30SiC and Ti₃SiC₂–30TiC (vol.-%),293 both TiC and SiC were found to lower grain boundary mobility in Ti₃SiC₂. At comparable grain sizes, all composites tested were weaker in flexure than the unreinforced Ti₃SiC₂ matrix, with the reduction in strength being the worst for the SiC composites, which was attributed to thermal expansion mismatches between the matrix and reinforcement phases. The composite samples, however, were exceptionally damage tolerant and thermal shock resistant.

Incorporating secondary phases into the MAX phase microstructure usually aims at strengthening the compound. However, the unique properties of the MAX phases can also be utilised as a binder phase for hard metals. For example, hard alloys for cutting tools and moulds, based on WC, are vital for the manufacturing industry. However, W is a rare metal. Therefore, there is a demand, particularly in Japan, for elements that can replace W. TiC is a carbide with better hardness than WC and is a good candidate in replacing WC in hard alloys. There have been some attempts for using TiC–Ni, cermets, as a substitute for W. This approach today, has not been widely accepted partially because Ni is also a relatively rare element and the properties of such cermets are not satisfactory yet.

Sun and co-workers synthesised Ti₃SiC₂ using TiC in the starting materials (e.g. Refs. 107, 112 and 114), and this motivated the in situ synthesis of TiC–Ti₃SiC₂ composites by adjusting the ratio of TiC in the starting powders. Better interface structures were expected in the composites synthesised via this route. TiC–Ti₃SiC₂ composites294,295 were synthesised by an in situ reactive PDS process. The strengths of the composites was increased by the incorporation of Ti₃SiC₂ and a maximum strength was achieved at ∼50 vol.-%Ti₃SiC₂. The grain size of the Ti₃SiC₂ phase was still relatively coarse (Fig. 15). Processes which produce composites with finer Ti₃SiC₂ grains may result in better materials. Processes such as MA followed by HP296 and a newly developed travelling zone sintering method297 were also employed for the TiC–Ti₃SiC₂ composites.

Figure 15.

Microstructure of TiC–55 vol.-%Ti₃SiC₂ composite294

Toughening SiC with Ti₃SiC₂ dispersions was also investigated.298 The addition of Ti₃SiC₂ to SiC led to improved fracture toughness, with the best combination of properties obtained to be a K _IC = 8·3 MPa m^1/2 and an HV = 17·6 GPa. Conversely, it is possible to reach a combination of hardness and fracture toughness by controlling the decomposition of Ti₃SiC₂ to various binary carbides.

Much less information on Ti₃AlC₂ based composites is available in the literature. Starting with Ti₃AlC₂ powders and microsized SiC and ZrO₂, a Ti₃AlC₂ composite was synthesised.299 Both the Ti₃AlC₂–SiC and Ti₃AlC₂–ZrO₂ composites were superior to the monolithic Ti₃AlC₂ in strength, fracture toughness, and microhardness.

MAX–oxide

Luo et al. 300 fabricated Al₂O₃–Ti₃SiC₂ composites with a full range of volume fractions, by attrition milling the Al₂O₃ and Ti₃SiC₂ powder mixtures followed by an spark plasma sintering process. The hardness decreased with increasing of Ti₃SiC₂ content, while the fracture toughness and the strength increased. It is interesting to note that the sample without Al₂O₃ showed the maximum flexural strength (673 MPa) and the maximum toughness (9·3 MPa m^1/2), which is relatively high for monolithic Ti₃SiC₂. A small amount of TiC was found in the microstructures and that was attributed to the reaction of carbon in the furnace with the Ti₃SiC₂. However, according to the experience of this review’s author, ball milling involved in the fabrication process may have a considerable effect on the formation of ancillary phases such as TiC. In the composites with Ti₃SiC₂ as matrix and up to 20 vol.-%Al₂O₃ as reinforcement,301 that was in situ synthesised from mixtures of reactant powders, the hardness increased considerably, while the dependency of strength and fracture toughness were generally in agreement with previous results300 though the absolute values are much lower.

Using a similar process,300 Pan et al. 302 fabricated Ti₃SiC₂–TZP composites with up to 50 vol.-%Ti₃SiC₂. While the Vickers hardness values and flexural strengths of the composites decreased with increasing Ti₃SiC₂ content, the fracture toughness increased. The maximum fracture toughness of ∼10 MPa m^1/2 was achieved for the composite with 50 vol.-%Ti₃SiC₂. The dielectric behaviour of Ti₃SiC₂–TZP composites was also reported.303

Composites of Ti₃AlC₂, with up to 20 vol.-%Al₂O₃, were fabricated by reactive sintering.304 The strengths, hardness and toughness values are enhanced by the incorporation of Al₂O₃. The strengthening and toughening effects were attributed to possible compressive residual stresses at the interfaces, a conjecture that may need more confirmation. When the composite with the optimal Al₂O₃ content (10 vol.-%) was tested at high temperatures, the strengthening effect caused by Al₂O₃ strongly depended on the mode of failure.305 The dielectric properties and the electromagnetic interference shielding effect of Ti₃SiC₂–Al₂O₃ composites were also reported.306,307

Combustion reaction of TiO₂–Al–C at ∼1173 K leads to Ti₃AlC₂–TiC–Al₂O₃ composites.308 Vickers hardness, fracture toughness, and flexural strength of the nearly dense sample synthesised from 3TiO₂–5Al–2C were measured to be 13·3 GPa, 5·8 MPa m^1/2 and 466 MPa respectively. This process, starting from low cost raw materials such as TiO₂, should be pursued more vigorously since it may prove to be a low cost method for the synthesis of MAX phases.

It is interesting to note a report on a three-dimensional printing process developed for the multiphase composite of Ti₃AlC₂–TiAl₃–Al₂O₃.309

MAX–boride

Cubic boron nitride (cBN) is the second hardest material next to diamond. Benko et al. 310 fabricated cBN–Ti₃SiC₂ composites and observed various phases such as TiB₂, SiC₄, TiC and SiC at the cBN/Ti₃SiC₂ interfaces.

TiB₂ also shows high hardness, modulus, low density, excellent chemical inertness, high electrical conductivity (5×10⁷–10×10⁷ Ω⁻¹ m⁻¹) and a coefficient of thermal expansion close to Ti₃AlC₂, that makes it a good reinforcement for Ti₃AlC₂. For example, the maximum compressive and flexural strengths of Ti₃AlC₂–TiB₂ composites were measured to be 2·2 GPa and 861 MPa respectively,311 in contrast to the respect values of monolithic Ti₃AlC₂ of 723 and 340 MPa. The strengthening effect was operational at least up to 1373 K. Contribution in electrical conductivity from TiB₂ was also observed.

MAX–metal

Via a powder metallurgy route, Zhang et al. 312,313 fabricated a Cu matrix composite reinforced by Ti₃SiC₂ particulates, aiming at improving the mechanical properties, without seriously decreasing the conductivity, for potential application as new electrofriction materials. The Cu–Ti₃SiC₂ composites showed superior mechanical properties over Cu–graphite composites. At filler contents of <20 vol.-%, the electrical conductivity for Cu–Ti₃SiC₂ composites was higher than that for Cu–graphite composites. At higher filler contents, however, the electrical conductivity for Cu–Ti₃SiC₂ composites was lower than that for Cu/graphite composites because of the presence of residual pores. Resembling Cu–graphite composites, the Cu–Ti₃SiC₂ was also found to be self-lubricant.

At high volume fractions of Ti₃SiC₂, however, densification of Cu–Ti₃SiC₂ composites became difficult.314 To solve this problem, Ti₃SiC₂ particulates were coated with a layer of Cu by an electroless plating method.315 The coated layer of Cu prevented direct contact and agglomeration of the Ti₃SiC₂ particulates. The composite reinforced with Cu coated Ti₃SiC₂ exhibited higher density and homogeneous distribution of Ti₃SiC₂ in the Cu matrix, and hence improved mechanical properties at ambient and elevated temperatures.314

Peng220 fabricated Cu–Ti₃AlC₂ composites and found that Ti₃AlC₂ increased the strength of Cu considerably. Unfortunately, the electrical conductivity was reduced by two orders of magnitude in comparison with the Cu matrix.220,316

Ti₂SnC was reported to show the highest electrical conductivity (14×10⁶ Ω⁻¹ m⁻¹),69 among the MAX phases (Table 1), and thus also employed as a reinforcement for strengthening Cu.317 ^– 319 The interfacial reaction between Cu and Ti₂SnC was observed to be controlled by the de-intercalation of Sn from Ti₂SnC to form Cu(Sn) solid solutions and TiCx.317 The Cu grain size was significantly decreased by incorporating the Ti₂SnC dispersion. Consequently, the yield strengths were enhanced by a factor of four over those of pure Cu.319 The most encouraging result is that the conductivity of the composite was 85·6% of pure Cu.319 The friction coefficients and wear rates of the Cu matrix composites were found to decrease significantly by the incorporation of the Ti₂SnC particles. For example, the wear rate of a Cu–10 vol.-%Ti₂SnC composite was only one-eighth of pure Cu.318 It thus appears that, to date, Ti₂SnC is the best possible MAX phases reinforcement for Cu as far as mechanical and electrical properties are concerned.

MAX composites for biomaterials

Hydroxyapatite has biological properties of great interest for osteo-implantation. However, its poor mechanical properties restrict its application. Considering the bioinertness of Ti₃SiC₂,320 Shi et al. 321 fabricated Ti₃SiC₂–hydroxyapatite composites with different Ti₃SiC₂ volume fractions. The composite’s flexural strengths and fracture toughness values increased with increasing Ti₃SiC₂ content, up to 252 MPa and 3·9 MPa m^1/2 respectively, at 50 vol.-%Ti₃SiC₂. On the other hand, the Vickers hardness decreased with Ti₃SiC₂ content, which favours bio-applications. The composites are machinable when the Ti₃SiC₂ content is >20 vol.-%.322

Carbide derived carbon (CDC)

Nanoporous CDC structures have been produced by selectively etching metal atoms from metal carbides.323 The fact that the porosity of CDC can be accurately controlled for specific postulated end uses has led to much effort to determine the effects of pore size, volume, and surface area on transport and adsorption properties. The tailored properties that can be achieved in CDC have facilitated advances in demanding applications such as hydrogen storage, supercapacitors and devices for treatment of sepsis extracorporeally (i.e. sorption of biomolecules).323

Beside processes based on binary carbides as the precursors, CDCs have also been synthesised through chlorination of carbon and nitrogen containing MAX phases at elevated temperatures.324 ^– 327 Gogotsi et al. 324 studied CDC formation from Ti₃SiC₂ powders and bulk samples as a function of synthesis temperature. Etching of Ti₃SiC₂ led to a larger pore volume (∼75%) compared with TiC (56·2%) or SiC (57·3%). Chlorination in a flow of pure Cl₂ for 3 h in a quartz tube furnace resulted in the extraction of Ti and Si from Ti₃SiC₂ leading to the formation of C by the reaction of Ti₃SiC₂+8Cl₂ (g) = SiCl₄ (g)+3TiCl₄ (g)+2C. Carbide derived carbons derived from Ti₃SiC₂ have a theoretical density of 0·55 g cm⁻³.

Carbide derived carbons synthesised starting with Ti₂AlC powders via chlorination have structures that are primarily amorphous when the chlorination temperature is low.325 Graphitic ribbons, as well as sharply bent graphitic structures, were observed at 1073 K. As the chlorination temperature was further increased to 1273 K, the width of the graphitic ribbon increased. No significant increase in graphitisation occurred between 1273 and 1473 K. Sorption measurements determined the presence of micropores (0·40–2·0 nm) after chlorination at 673 K; chlorination at 1073 K or higher resulted in both micro- and mesopores (0·35 to >7 nm).325

A systematic characterisation was carried out for both the micro- and mesoporosity, that developed when Ti₃SiC₂, Ti₃AlC₂, Ti₂AlC, Ti₂AlC_0·5N_0·5, Ta₂AlC, as well as two binary carbides of Ta₂C and TaC were chlorinated at different temperatures.323 The CDCs’ porosity depended on the initial carbide’s structure and chemistry. Carbides with anisotropic structures, such as the MAX phases, resulted in larger mesopore volumes as compared to more isotropic carbides such as TaC. Furthermore, the carbon pore structure can collapse, resulting in low pore volumes, for materials with very low molar fractions of carbon, as was the case with Ti₂AlC_0·5N_0·5.

Applications

Reviewing the properties of the MAX phases in 2000, Barsoum1 listed potential applications as including substitution for machinable ceramics, kiln furniture, wear and corrosion protection, heat exchangers, applications where rotating parts are used, low friction applications based on basal plane lubricity.

El-Raghy and Barsoum founded the company 3-ONE-2, LLC, apparently after the 312 phases, in 2001 as a joint venture with Kanthal AB (a Sandvik subsidiary). Since then, Maxthal 312 (Ti₃SiC₂) and Maxthal 211 (Ti₂AlC) have been commercialised in the form of powders and fully dense parts. Currently the full scale pilot plant is capable of producing batches of >100 kg of −325 mesh powders of Ti₃SiC₂ and Ti₂AlC.

Beside the commercialisation of the Maxthal powders, heating elements made of Maxthal 211 were developed (Fig. 16a ), and this was extended to the novel high temperature reactors for in situ X-ray absorption spectroscopy measurements in fluorescence (Fig. 16b ).328

Figure 16.

Some examples of MAX phases applications

Maxthal 211 was also shown to be superior to steel when used as gas burner nozzles, in sulphur containing natural gas corrosive burning environments. Figure 16c compares nozzles made of Maxthal 211 and 353MA steel after one year in a furnace at 1473 K. The α-Al₃O₂ scale that forms on the Ti₂AlC protected the material and allowed an increase of the process temperature up to 1673 K. Because of the excellent machinability, threading enabled the MAX phases nozzles to directly replace the steel ones.

A project funded by the Office of Navy Research led to the successful development of MAX based materials for foil bearings, having low friction and wear from room temperature to 823 K. Figure 16d shows a Ta₂AlC/Ag shaft and a superalloy foil after testing for 10 000 cycles in a rig. The shaft was rotated at 50 000 rev min⁻¹.329 Honeywell Inc. is further developing this concept.

3-ONE-2 and Hilti developed tooling for the dry drilling of concrete, consisting of diamonds in MAX 312 segments, which were then brazed to steel. The performance of the MAX phase segments was reported to be far superior to that of current diamond/Co segments (Fig. 16e ). With further improvements in design to overcome the problem of smearing of concrete powders due to high temperatures and the inadequate toughness of the segments, this material may be close to market.

The success of slip casting of MAX phases has extended the potential application of these materials. For example, three-dimensional shaping and sintering to final size of fully dense and hollow objects has been achieved (Fig. 16f ).

Owing to their good electrical conductivity and tribological properties, in addition to the acceptable mechanical properties, MAX phases such as Ti₃SiC₂ and Ti₃AlC₂ have been shown to perform better than carbon based pantographs for electric trains. Currently, a few leading projects are running in China, aiming for application on the high speed railway under construction.

In addition to the applications mentioned above, a number of other potential applications have been mentioned in previous sections. These include: electrodes, exhaust gas filters for automobiles, free-cutting elements, microelectronics, biomaterials, damping materials (high stiffness and up to high temperatures), corrosion resistant materials, surface coatings, defence applications, such as armour, nuclear applications, low dimensional materials, and substrates for CVD diamond.

Development of key technologies of MAX phases, such as pressureless sintering, surface coating by PVD and MAX based composites, is considered most likely to lead to more breakthroughs in the application of MAX phases in the near future.

Future directions

The most urgent issue in MAX phase research is to develop viable commercial applications for these materials. With the preliminary and potential application fields discussed above, market penetration by MAX phase materials will create more stimuli to carry out more research and development on this family of layered metallic ceramics. Research on MAX phases is currently funded by agencies, science foundations and defence industries in the USA, China, Europe and Australia. However, funding for research on MAX phases from governmental agencies in Japan is almost zero. The likelihood of the Japanese government funding such research remains small until a clear market scenario is available.

Another important issue is the search for new MAX phases. First, to better understand the MAX phases themselves, both those known to exist and those theoretically predicted to be thermodynamically and mechanically stable. To do so, thorough experimental investigations into the properties are needed. The second priority is to discover more MAX phases. It is encouraging that in the past few years many MAX phases, particularly 413 and 312 phases, have been discovered. More 312 or 211 MAX phases, plus their solid solutions, which are numerous, are awaiting exploration.

Cost is also a challenge. However, with the diverse selection of the M, A and X elements, and the increasing diversity in starting reactants such as binary M–X, M–X, A–X compounds and even oxides, MAX phases with reasonable cost are not out of reach. In addition, processes such as pressureless sintering provide decreased cost, with the possibility of fabricating MAX phases in larger quantity and complex geometries. This has already been shown to be possible for some products as demonstrated by the 3-ONE-2 and Kanthal joint venture.

Footnotes

Acknowledgements

Contributions from the author’s colleagues Dr H. Hashimoto, Dr T. Abe, and postdoctoral research fellows Dr Z. F. Zhang, Dr S. L. Yang, Dr Y. Zou, Dr Y. L. Du, Dr W. B. Tian to the studies on the MAX phases are appreciated. The inspiration provided by collaboration with Dr M.W. Barsoum is also acknowledged; his advice and discussion have been invaluable. Dr M. W. Barsoum and Dr Y. C. Zhou are acknowledged for their critical reading of and comments on this manuscript. The author would also like to thank Nihon Dennetsu Co. Ltd for financial support.

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Progress in research and development on MAX phases: a family of layered ternary compounds

Abstract

Introduction

MAX phases reported in literature as having been experimentally synthesised, proposed and/or theoretically predicted, and some fundamental properties*: solid solutions, except those without end members reported, are not listed; polymorphs calculated as unstable are not listed

Synthesis

Synthesis of bulk MAX phases

Synthesis of MAX phase films

Synthesis of MAX phase powders

Properties

Polymorphism

Phase stability

Electrical, thermal and elastic properties

Hardness, machinability, deformation and fracture

Fatigue and creep

Oxidation

Parabolic oxidation rate constants of some typical MAX phases

Corrosion

Friction and wear

Other surface phenomena

Materials and processes based on MAX phases

Surface technology

Surface modification

Diamond on Ti3SiC2

Joining

Composites

MAX–MX and MAX–AX

MAX–oxide

MAX–boride

MAX–metal

MAX composites for biomaterials

Carbide derived carbon (CDC)

Applications

Future directions

Footnotes

Acknowledgements

References

Diamond on Ti₃SiC₂