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Plasma Nitriding of Al–Cu and Al–Cu–Mg Alloys at Various Holding Temperature

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Tatsuhiko Aizawa, Hiroki Nakata and Takeshi Nasu

Submitted: 10 September 2025 Reviewed: 25 November 2025 Published: 04 March 2026

DOI: 10.5772/intechopen.1014134

Aluminium Alloys - Synthesis, Properties, and Applications IntechOpen
Aluminium Alloys - Synthesis, Properties, and Applications Edited by Shashanka Rajendrachari

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Aluminium Alloys - Synthesis, Properties, and Applications [Working Title]

Shashanka Rajendrachari

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Abstract

A plasma nitriding has been highlighted as a surface treatment to harden aluminum alloys and improve their wear resistance. Various processing approaches were reviewed to describe the nitriding processes at relatively high temperatures. Duralumin and super-duralumin alloys, such as Al–Cu and Al–Cu–Mg alloys, were nitrided by DC- and RF-DC plasmas at holding temperatures ranging from 723 K to 623 K. An AlN (Aluminum Nitride)-rich layer formation characterized the nitriding behavior of Al–Cu alloys at 673 K. The intermetallic precipitate of Al2Cu facilitated this fast-rate AlN-layer formation as a catalyst. This process was affected at 723 K by Mg diffusion and surface oxidation in the case of plasma nitriding of Al–Cu–Mg alloys. The plasma nitriding at 623 K was free from this deterioration, allowing the AlN-rich layer to form in a manner similar to the nitriding of Al–Cu alloys. This layer exhibited a hardness greater than 1,000 HV, which is 10 times higher than the original hardness of Al–Cu–Mg alloys. A heatsink was plasma nitrided at 623 K to demonstrate that aluminum alloy parts could be homogeneously nitrided to form the AlN-rich layer. The plasma nitriding at temperatures lower than 600 K was investigated to explore the possibility of nitrogen supersaturation in aluminum alloys, as well as the accommodated nitriding process via intermetallic precipitates other than Al2Cu.

Keywords

  • plasma nitriding
  • intermetallic precipitates
  • AlN-rich layer formation
  • surface hardness
  • holding temperature effect
  • duralumin alloys

1. Introduction

Light-weight metals and alloys have been highlighted as structural materials for electric vehicles, rail-road trains, and tracks to save energy consumption and reduce environmental burdens [1]. Among them, aluminum and aluminum alloys are worth using as sustainable materials due to their recyclability with minimum energy consumption [2]. In the life cycle assessment of aluminum and aluminum alloy products, their reusability must be improved by increasing wear resistance via surface treatment [3]. Anodizing process [4] was invented to synthesize an alumina layer onto the surface of various aluminum products. In addition to its residual porosities, its low toughness has been identified as a fatal drawback that shortens the life time of aluminum alloy products.

Instead of the oxidation process, various nitriding treatments have been proposed to harden steels [5]; e.g., as listed in Table 1, the gas nitriding using ammonia [6], liquid nitriding using chloride solutions [7], and plasma nitriding using a nitrogen–hydrogen mixture gas [8]. Among these, plasma nitriding provides an environmental-friendly method to improve surface hardness and tribological performance, as compared in Table 1.

Nitriding Source TemperatureK Durationks HardnessHV Thicknessμm Work
Salt bathnitriding NaCNKCN 833–853 180–360 1000 100–300 Alloysteel
Gasnitriding NH3 733–853 90–360 700–1,200 100–800 Highalloysteels
Plasmanitriding N2 + H2 623–863 3.6–108 600–1,300 50–90 Steels

Table 1.

Comparison of three nitriding processes for surface treatment of steels.

In the literature, various methods have also been proposed for the surface treatment of aluminum and aluminum alloys. Table 2 compares three plasma nitriding processes used to form a nitrided surface layer on aluminum and aluminum alloys. Relatively high holding temperatures and longer nitriding durations were common to these three processes. This fatal drawback might stem from the very low diffusion coefficient of nitrogen solutes through AlN [8].

Method Performance Merits Demerits REF
Ion implantation Implantation energy(>10KeⅤ)Lattice distension of AlN No need forpresputtering High energyconsumptionDifficulty inapplication tocomplex-shaped partsSurface damage [9, 10]
Plasma immersion Ion Implantation 10–15 μm thick film723–823 KMore than 18 ks4.5 μm by 43.2 ks Short processingtimeThick AlN layer Presputtering(36ks)Surface damageHigh energyconsumptionDifficulty inapplication toautomotive parts [11, 12]
Plasma nitriding Several microns in thicknessMore than 723 KLong than 28.8 ks Low energyconsumptionLow gasconsumptionDirect application to industry Pre-sputtering by ArLong nitriding timeRelatively highholding temperature [13, 14]

Table 2.

Comparison of three plasma nitriding processes for aluminum and aluminum alloys.

In the literature, many efforts have also been made to modify the normal nitriding process for the synthesis of a thick AlN layer at a lower holding temperature and for a shorter duration. Three processes have been cited in the literature to describe this improvement in nitriding characteristics. As listed in Table 3, DC (direct current)-plasma nitriding, with the use of screens, played a role in nitriding various steels; however, little data has been reported on further improvement over the normal plasma nitriding described in Table 2 [15]. RF (Radio Frequency)-plasma nitriding succeeded in lowering the holding temperature, and a very thin AlN-layer was synthesized using this approach [16]. With the assistance of the electron beam, plasma nitriding also succeeded in the synthesis of a thick AlN-layer at a relatively lower holding temperature and for a shorter duration [17].

Table 3.

Comparison of three plasma nitriding processes for the advancement of nitriding characteristics.

In the present chapter, the nitriding behavior is reconsidered to search for an efficient plasma nitriding treatment at a lower holding temperature for a shorter nitriding duration. DC-plasma and RF–DC plasma nitriding processes are utilized to describe the nitriding behavior of Al–Cu and Al–Cu–Mg alloys in the 2000-series of aluminum alloys. The role of intermetallic precipitate (IMP) is discussed as the main mechanism driving the formation of a thick AlN-layer in both the DC- and RF–DC plasma nitriding processes. Microstructure analyses are performed to experimentally prove this Al2Cu-accommodated formation of the AlN layer via the reactions by Al2Cu + 2 N (diffusing nitrogen solute atoms) → 2AlN + Cu, and, Cu + 2Al (in matrix) → Al2Cu. Al2Cu works as a catalyst to drive the overall reaction by Al (in matrix) + N (diffusing nitrogen solute atom) → AlN. This reaction takes place at the interfacial boundaries between AlN and Al2Cu synthesizing a thick composite layer of AlN and Al2Cu within the Al–Cu and Al–Cu–Mg alloy matrices.

In addition to this positive role of IMP, the alloying element in these alloys influences the nitriding behavior. In the case where the Al–Cu–Mg alloys are RF–DC plasma nitrided at 723 K, the surface diffusion of Mg and its oxidation hinder the nitriding process. This deterioration is suppressed by lowering the temperature to 623 K to make AlN-layer formation active in nitriding the Al–Cu–Mg alloys. A duralumin heat sink is successfully nitrided at 623 K, demonstrating that aluminum alloy parts are hardened by the formation of an AlN-layer.

Two RF–DC plasma nitriding approaches are discussed for 5,000, 6,000, and 700 series aluminum alloys, other than Al–Cu and Al–Cu-Mg alloys, at low holding temperatures. A massive nitrogen supersaturation provides a candidate treatment to build up the AlN-layer using the in-situ granular refinement and fine grain- and cluster-boundary network for nitrogen diffusion. IMP, other than Al2Cu works as a catalyst to drive the reaction by Al + N → AlN with less dependency on the chemical compositions of aluminum alloys.

2. Methods and materials

Two plasma nitriding systems were introduced to form AlN-rich surface layer on Al–Cu and Al–Cu-Mg alloys; e.g., DC plasma and RF–DC plasma nitriding systems. The standardized aluminum alloys, such as AA2011 or AA2014 were utilized as specimens, as well as the home-made alloys.

2.1 DC-plasma nitriding system

A DC-plasma nitriding system without the use of screens [15] was employed to form an AlN-rich layer on the Al–Cu and Al–Cu-Mg alloy specimens, which were prepared by their own melting and casting processes. Their chemical components are listed in Table 4. As shown in Figure 1a, the nitrogen – hydrogen plasma was ignited by glow discharge under the applied DC bias voltage. The specimen was set up on the DC-biased plate, below which the electrical heater was located. A thermocouple was used to measure the specimen temperature, using a dummy specimen [18].

Figure 1.

Schematic views of two plasma nitriding systems. a) DC-plasma nitriding system, and b) RF–DC plasma nitriding system.

Cu Mg AI
AI-Cu alloy 6.0 0.0 Bal.
AI-Cu-Mg alloy 6.0 0.5 Bal.

Table 4.

Chemical compositions of Al–Cu and Al–Cu–Mg alloy specimens.

A typical nitriding condition is listed in Table 5. To be discussed later, nitrogen and hydrogen gas pressures are fixed to be high enough to sustain the yield of NH-radicals in DC-plasma nitriding [19, 20]. In reference to ammonia gas, which is usually employed in the so-called ion- and radical-nitriding processes in industries [21, 22], the gas flow rate in the mixture gas was also fixed to be 25% for nitrogen gas and 75% for hydrogen gas.

Process Parameters
Pre-sputtering DC (-500 V)Pressure (133 Pa)Temperature (723K)Duration (1.8 ks), Carrier gas (N2 only)
Nitriding DC (-500 V)Pressure (533 Pa)Temperature (673K)Duration (7.2, 14.4, 21.6, 28.8 ks)Carrier gas (N2 + H2)Partial pressure ratio (N2:H2 = 1 : 3)

Table 5.

DC-plasma nitriding conditions.

2.2 RF–DC plasma nitriding system

A RF–DC plasma nitriding system using the hollow cathode was employed to describe the effect of holding temperature on nitriding behavior, as shown in Figure 1b [23]. A nitrogen–hydrogen plasma was ignited only in the hollow so that the nitrogen ion and NH-radical densities were enhanced to increase the nitrogen solute flux into AA2011 or AA2014 alloy matrices. Their chemical compositions are listed in Table 6. A typical plasma nitriding condition is summarized in Table 7. Compared to the DC-plasma nitriding conditions in Table 5, lower gas pressure was utilized, around 75 Pa. The gas flow rate was determined by the in-situ plasma diagnosis to be 80% for nitrogen gas and 20% for hydrogen gas in mixture [24].

Si Fe Cu Mn Mg Cr Zn Ti Ti+Zr AI
A2011 0.15 0.23 5.42 0.00 0.00 0.00 0.01 0.03 ------ Bal.
A2014 0.8 0.3 4.5 0.72 0.48 0.02 0.08 0.02 0.02 Bal.

Table 6.

Chemical compositions of AA2011 and AA2015 aluminum alloy specimens.

Process Parameters
Pre-sputtering DC(-500 V), RF (0 V)Pressure (30 Pa)Temperature (623K, 673K, 723 K),Duration (1.8 ks), Carrier gas (N2 only)
Nitriding DC (-500 V), RF (250 V)Pressure (75 Pa)Temperature (623K, 673K, 723 K),Carrier gas (N2 + H2)Partial pressure ratio (N2:H2 = 4:1)

Table 7.

RF–DC plasma nitriding conditions.

2.3 Mechanical characterization

Micro-hardness testing was employed to analyze the mechanical response of DC- and RF–DC plasma nitrided aluminum alloy specimens. The hardness depth profile was measured to describe the advancement of the nitrided layer into the depth of the matrix. The surface hardness, by varying the applied load, was also measured to analyze the hardening behavior due to AlN-formation during nitriding.

2.4 Microstructure characterization

SEM (scanning electron microscopy) – EDS (energy dispersive X-ray spectroscopy), as well as TEM (transmission electron microscopy), were utilized for microstructure characterization.

XRD (X-ray diffraction) and GIXD (glancing incident X-ray diffraction) were utilized to identify the synthesized phases during nitriding. XPS (X-ray photoelectron spectroscopy) was also used to analyze the synthesized compounds on the nitrided specimen surface, to describe the depth profile of related elements in the nitriding process, and to measure the binding state of compounds at the in situ-formed surface layer of the nitrided specimen.

3. DC-plasma nitriding of Al–Cu alloys at 673 K

DC-plasma nitriding was first employed to describe the nitriding behavior of Al–Cu alloys at 673 K. The nitriding duration was varied to examine the nitriding transients. An AlN-layer was formed as a hardened surface layer by DC-plasma nitriding for 28.8 ks. The nucleation and growth mechanism of the AlN-layer was experimentally analyzed to discuss the role of Al2Cu precipitates in the AlN-layer formation.

3.1 Materials characterization

GIXD (grazing incident X-ray diffraction) was utilized to describe the nitriding process of Al–Cu alloys with increasing nitriding duration. As shown in Figure 2, the XRD profile changed significantly with duration. Peaks for aluminum nitride were detected with small intensity in the nitrided sample at 7.2 ks; the main peaks come from aluminum. This implies that AlN only nucleates along the grain boundaries. With increasing nitriding time, the peak intensity ratio of AlN to Al monotonically increases. At 14.4 ks, AlN and Al2Cu grew to become the main components together with the aluminum matrix. After 14.4 ks, these aluminum peaks became relatively small, indicating that the nitrided layer consists of AlN and Al2Cu.

Figure 2.

Variation of GIXD diagrams for DC-plasma nitrided Al–Cu alloy at 673 k with increasing the nitriding duration.

Figure 3a depicts the surface of a DC-plasma nitrided Al–Cu alloy specimen at 673 K for 28.8 ks. The original aluminum alloy surface consists of a nitrided layer without the Al–Cu alloy matrix. This corresponds to the GIXD profile in Figure 2, where only AlN and Al2Cu peaks are detected, without aluminum peaks. As shown in Figure 3b, the cross-section of the nitrided layer has a columnar structure. This thick nitrided layer has a columnar structure, and the length of this columnar structure is approximately equal to its thickness. The diameter of each column is around 20–30 μm, equivalent to the original grain size of the Al–Cu alloy. Through GIXD and EDX analyses, this column is composed only of AlN and Al2Cu. This suggests that AlN nucleates along the grain boundaries and grows into the Al–Cu alloy matrix via the formation of a columnar composite of AlN and Al2Cu.

Figure 3.

SEM image on the DC-plasma nitrided Al–Cu alloys at 673 k for 28.8 ks. a) Surface of nitrided layer, and b) cross-section of nitrided layer.

The large yield of AlN in Figures 2 and 3 is a proof that nitrogen atoms are supplied to the nitriding front end with sufficient nitrogen content to sustain the nitriding reaction. The nitrogen diffusion coefficient becomes low in AlN. In addition, no nitrogen is solvable in Al2Cu. Therefore, interfacial boundaries between the formed AlN and Al2Cu become an effective diffusion path for nitrogen in this nitriding process. Let us analyze this interfacial boundary between AlN and Al2Cu in the following.

A crystallographic analysis was conducted to describe the coherency between adjacent Al2Cu and AlN. As depicted in Figure 4a, numerous interfacial zones are observed among AlN columns, where AlN and Al2Cu are present in adjacency. AlN has a hexagonal structure with a = 0.497 nm and b = 0.497 nm, as shown in Figure 4b. Meanwhile, Al2Cu or θ-phase in the phase diagram of Al–Cu base alloys, has a tetragonal structure with a = 0.606 nm and c = 0.487 nm. For the estimation of the crystallographic coherency between AlN and Al2Cu, their d-spacing distances were calculated. As demonstrated in Figure 4c, in the case of AlN, the d-spacing between {-101} and {-11-1} planes is 0.237 nm, and the d-spacing between {121} and {211} planes is 0.236 nm in the case of Al2Cu. Both values are in fairly good agreement. This confirms that AlN is synthesized in coherency with Al2Cu. AlN is formed by the reaction between Al2Cu and diffusing nitrogen solute through the interfacial boundaries; e.g., Al2Cu + 2 N → 2AlN + Cu. Hence, metallic copper must be present at the interfacial boundaries. Let us perform XPS analysis on the copper binding state at the B-zone.

Figure 4.

Coherency between adjacent AlN and Al2Cu. a) Interfacial boundary B among AlN columns, b) crystallographic structure of AlN in the A-zone, and c) crystallographic structure of AlN and Al2Cu in the B-zone.

As shown in Figure 5, metallic copper coexists with Al2Cu; the AlN formation reaction advances in the interfacial zones via Al2Cu + 2 N → 2AlN + Cu. Remembering that the initial content of Cu is only 6 mass%, a significant amount of Al2Cu is reproduced in the nitrided layer during the nitriding process. Residual Al2Cu and Cu in the nitrided columnar structure in Figure 5 also reveal that the solid-state reaction, represented by Al2Cu  + 2 N balu 2AlN + Cu and 2Al (in matrix) + Cu balu Al2Cu takes place during this nitriding process for the formation of AlN and the reproduction of Al2Cu, respectively. This solid-state reaction becomes a driving process, utilizing Al2Cu as a catalytic material in plasma nitriding.

Figure 5.

XPS analysis of the copper binding state at the b-zone in Figure 4a.

3.2 Mechanical characterization

The micro-Vickers hardness testing was used to describe the transients of the hardness depth profile with increasing nitriding duration. Figure 6 depicts the advancement of hardness depth (h) with increasing nitriding time. It grows linearly with E; e.g., dh/dτ becomes nearly constant. As stated in [25, 26], E2 is theoretically or experimentally proportional to the nitriding duration (τ) if the nitriding behavior is mainly governed by nitrogen diffusion. Figure 6 implies that the fast-rate nitriding process is never simply governed by nitrogen diffusion when it is accommodated by the catalytic reaction of Al2Cu precipitates.

Figure 6.

Transients of hardness depth profile with increasing nitriding time.

The thick nitrided layer is a composite material of AlN with Al2Cu, the volume fraction of which is estimated to be 10 vol% in total. The mixture rule of hardness is used to estimate its hardness from 1,400 HV for pure AlN and 250–350 HV for Al2Cu [27]. Its measured hardness, 1200 HV in Figure 7, is in good agreement with the estimated value of 1200–1300 HV.

Figure 7.

Emissive light spectrum of the ignited nitrogen–hydrogen plasmas in the narrow wavelength range from 320 nm to 350 nm.

4. RF–DC plasma nitriding of Al–Cu alloys at 673 K

Al–Cu alloys, or AA2011 alloys, were employed to describe the plasma nitriding behavior at 673 K. In a similar manner to DC-plasma at 673 K, the AlN-layer was formed by the catalytic role of Al2Cu precipitates.

4.1 Plasma diagnosis

The emissive light spectroscopy (PMA-11, Hamamatsu-Photonics Co. Ltd.; Hamamatsu, Japan), as well as the Langmuir probe system (ALP System, Impedance Co. Ltd., Irland) were instrumented to the present RF–DC plasma nitriding system to make an in situ quantitative diagnosis of the generated nitrogen–hydrogen plasmas. In particular, a high-resolution spectroscopic analyzer with a resolution of 0.1 nm was utilized to detect the NH radicals in the specified wave length range. Furthermore, SPECAIR [28] was used as a simulator of emissive light spectra to theoretically predict the synthesis of NH-radicals in plasma nitriding under the mixture of nitrogen and hydrogen gases.

In this quantitative diagnosis of plasmas, the hydrogen-to-nitrogen gas flow rate was varied as a parameter with significant influence on NH-radical synthesis, as noted in [29, 30]. First, the emissive light spectra from nitrogen-hydrogen plasma were detected and analyzed using high-resolution spectroscopy under the standard experimental conditions outlined in Table 7. When scanning the measured spectra across a wider wave length range, the activated nitrogen atom (N* or N(I)), nitrogen ions (N+ or N(II)), and N2+ (N(III)) were detected to have significant peaks, along with the ionized nitrogen molecule (N2+). This synthesis of nitrogen ions is attributed to a series of reactions originating from the parent species N2+; e.g., N2+ → N* + N+, N+ → N2+ + e, and N+ + e → N*. Next, the wave length range was narrowed to 290–350 nm. The NH radicals were detected at λ = 336 nm with relatively high intensity, comparable to N2+ peak at λ = 337.1 nm, as depicted in Figure 7. The ratio of NH-radical peak intensity to N2+ peak intensity increases with decreasing hydrogen content; the hydrogen gas flow rate relative to nitrogen gas is expected to be a key parameter in nitriding, which will be discussed later.

In parallel with this experimental diagnosis, the program SPECLAIR was employed to simulate the NH radical spectra, as referenced in Refs. [31, 32]. Among several input parameters for this simulation, the electron temperature (Te) and its density (Ne) were directly measured using the Langmuir probe [33]; e.g., Te = 5 eV and Ne = 4 x 1016 m−3. The experimentally analyzed spectra of NH radicals generated in the nitrogen–hydrogen gas mixture with a ratio of four to one were in good agreement with the theoretically estimated spectra by SPECAIR, as seen in Figure 8. This assures that a significant intensity of NH radicals is generated under these plasma conditions with lower hydrogen content. The high electron density reveals that nitrogen ionization and NH-radical formation are enhanced under this plasma nitriding condition. That is, the experimental conditions in Table 7 were suitable for generating plasmas with the flux of N*, N+, N2+, N2+, and NH-radicals.

Figure 8.

Comparison of the emissive light spectra between the experimental measurement and the theoretical simulation.

4.2 Plasma nitriding of AA2011 at 723 K and 673 K

AA2011 Al–Cu alloy specimens were employed for RF–DC plasma nitriding at 723 K and 673 K. Figure 9 depicts the XRD diagram of the nitrided AA2011 disc specimen at 723 K. In addition to the aluminum peaks with the first and second highest intensities, respectively, at 2θ = 38.5° and 44.8°, typical triplet peaks for aluminum nitride (AlN) were detected at 2θ = 33.4°, 36.1°, and 38.0°, respectively. This assures that the nitrided layer consists of a composite of synthesized AlN phase and AA2011 matrix. Comparing this XRD diagram to that of DC-plasma nitrided Al–Cu alloy for 14.4 ks, both are in fairly good agreement with each other. This implies that nearly the same nitriding process advances into the depth of the Al–Cu alloy matrix in both plasma nitriding treatments.

Figure 9.

XRD diagram of RF–DC plasma-nitrided AA2011 specimen at 723 K for 14.4 ks.

The plasma-nitrided AA2011 disc specimen, under the experimental conditions outlined in Table 7, is shown in Figure 10a. The specimen surface was homogeneously nitrided, resulting in a black coloration. Figure 10b depicts its microstructure. The black zone corresponds to the nitrided area, while the white zone represents the original matrix. The area fraction of the nitrided zones relative to the whole surface area is 80%, as shown in Figure 10b. Subsequently, its volume fraction (f) is estimated to be approximately 60%. The mechanical properties of this composite microstructure are evaluated through micro-hardness testing in the following sections.

Figure 10.

Optical-microscopy and SEM observation of the RF–DC plasma nitrided AA2011 specimen at 723 K for 14.4 ks.

Most of the copper in AA2011 or Al–Cu binary alloy is bound in the form of Al2Cu as a precipitate. This Al2Cu first precipitates during hot extrusion and exists in the vicinity of the grain boundary. This original precipitate still coexists with AlN, as discussed in Refs. [18–20]. This Al2Cu-precipitate also works as a catalyst in the RF–DC plasma nitriding to accommodate the fast-rate nitriding by a series of solid-phase reactions: Al2Cu + 2 N → 2AlN + Cu, and Cu + 2Al → Al2Cu. Then, an AlN-cluster is expected to nucleate at the grain boundaries, including Al2Cu precipitates. To describe this AlN-cluster formation, the presputtering temperature and the holding temperature were reduced to 673 K in Table 7 to prepare the nitrided AA2011 specimens in the early stage of nitriding by controlling the nitriding duration for microstructural analysis.

Figure 11 shows the SEM image of nitrided AA2011 specimen surfaces at 673 K by varying the nitriding duration (τ). When τ = 600 s, the AlN embryos are formed at the grain boundaries, as seen in Figure 11a. This proves that both the nitrogen solute isolating from NH radicals and the nitrogen ions react with Al2Cu precipitates and nucleate the AlN embryos at the grain boundaries through the chemical reactions by Al2Cu + 2 N → 2AlN + Cu and Cu + 2Al → Al2Cu. When τ = 3.6 ks, these embryos coalesce with each other and form a thin aln-layer on the granular surface of AA2011 specimen, as seen in Figure 10b. This observation in Figure 11 reveals that the nitriding process of Al–Cu alloys advances by the nucleation and growth mechanism of AlN-clusters under the intense nitrogen flux from the surface during the RF–DC plasma nitriding.

Figure 11.

AlN nucleation and growth processes in the RF–DC plasma nitriding of AA2011 at 673 K. a) SEM image on the AlN nucleation along the grain boundaries at τ = 600 s, and b) SEM image on the AlN growth at τ = 3.6 ks.

4.3 Mechanical characterization of RF–DC plasma nitrided AA2011 at 723 K

Variation of surface hardness with increasing weight (W) of the indenter is shown in Figure 12. In the case where W = 0.1 N or 10 gf, the surface hardness was measured to be 800 H. This hardness is eight times greater than the original hardness of AA2011 alloy. With increasing weight of the Vickers indenter, the measured surface hardness decreased. According to [34], the matrix hardness influences the measured hardness of materials since the matrix, six times deeper than the indentation depth, is forced into elasto-plastic deformation during indentation for hardness measurement. From this gradual decrease in surface hardness with W, the equivalent depth of the nitrided layer was estimated to be 15 μm when nitriding at 723 K for 14.4 ks.

Figure 12.

Variation of surface hardness with increasing the applied load in micro-vickers testing.

The nitrided layer becomes a composite of AlN-zones and an Al–Cu alloy matrix, as shown in Figure 10b. Assuming that the AlN-zones are homogeneously distributed in the entire nitrided layer by the volume fraction (f), as seen in Figure 10b, the linear law in homogenization theory [35] provides that the theoretical hardness (HT) is estimated by HT = (1-f) x Hmatrix + f x HAlN, where Hmatrix is the hardness of AA2011 matrix, and HAlN is the hardness of AlN. Since Hmatrix = 100 HV, HAlN = 1400 HV, and f = 0.6, HT is estimated to be 860 HV. This HT is in good agreement with the measured surface hardness, which ranges from 800 to 900 HV. This agreement confirms that the nitrided layer should be a composite of AlN and an Al–Cu alloy matrix, with a thickness of 15 μm.

5. RF–DC plasma nitriding of Al–Cu–Mg alloys at 723 K and 623 K

Most of the alloying elements in the aluminum alloys with lower melting temperatures have an influence on the nitriding behavior. AA2014 or Al–Cu–Mg alloy, was employed to describe this influence of Mg on the nitriding of the aluminum matrix. The holding temperature (TH) was varied to describe the thermal activation effect on the nitriding behavior. Magnesium diffusion is easily activated at TH = 723 K; however, it is retarded at TH = 623 K.

5.1 RF–DC plasma nitriding of Al–Cu–Mg alloys at 723 K

AA2014 specimen was also prepared for plasma nitriding at 723 K for 14.4 ks. The nitrided layer thickness was too thin, only to have the surface hardness from 100 to 120 HV. XPS was utilized to investigate this thin film formed on the AA2014 surface.

The chemical binding state of N1s in the vicinity of the nitrided surface is depicted in Figure 13a. The binding energy at the highest peak is 397 eV, corresponding to 397.3 eV for the binding energy of N1s in AlN. Figure 13b depicts the chemical binding state of Al2p. Since the binding energy of Al2p in metallic aluminum, Al2O3 and AlN is 72.7 eV, 74.1 eV, and 74.0 eV, respectively, the aluminum in this thin film is bound in AlN and Al2O3. This XPS analysis reveals that AlN-formation is significantly hindered during plasma nitriding at 723 K by the surface oxidation reaction.

Figure 13.

XPS analysis of the surface of RF–DC plasma-nitrided AA2014 specimen at 723 K. a) chemical binding state of nitrogen, and b) chemical binding state of aluminum.

XPS, with an etching rate of 1–2 nm/min, was utilized to perform chemical analysis from the surface to the depth. At each etching time, Al2p, Cu2p, Mg2s, O1s, and N1s diagrams were detected to calculate the integrated area of each peak profile and deduce the concentration of each of the five elements. Figure 14 depicts the variation of Al, Mg, Cu, N, and O concentrations with the etching time in the vicinity of AA2014 alloy surface after plasma nitriding at 723 K under the conditions outlined in Table 7. The nitrogen and copper contents were detected at sufficient concentrations to synthesize AlN through the reaction: Al2Cu + 2 N → 2AlN + Cu, and Cu + 2Al → Al2Cu. The magnesium content reaches 3–5 at%, which is much higher than the magnesium composition listed in Table 6. The high oxygen content detected in Figure 14 poses a risk of deteriorating the nitriding reaction due to surface oxidation, as seen in Figure 13b. As suggested in [21], magnesium solutes easily diffuse onto the surface through the grain boundaries, blocking the reaction between aluminum and diffusing nitrogen, and assisting the oxidation reaction between aluminum/magnesium and oxygen.

Figure 14.

Variation of the related elements (Al, Cu, Mg, O, and N) to nitriding behavior of AA2014 with the depth from the surface.

Although AA2011 or Al–Cu binary duralumin, was successfully nitrided under the same conditions as those in Table 7, the nitrogen diffusion and its reaction to synthesize AlN were completely retarded by magnesium concentration and surface oxidation during plasma nitriding at 723 K. A higher concentration of magnesium than the original content of 0.5 mass% might be attributed to the intense diffusion of magnesium solute from the depth of AA2014 matrix to the surface during nitriding at 723 K (or 450°C). The magnesium solute diffuses to the surface and reacts with oxygen to form magnesium oxides at this holding temperature [36].

5.2 RF–DC plasma nitriding of Al–Cu–Mg alloys at 623 K

Considering that this retardation by Mg-concentration at the surface should be only thermally controlled, the decrease of holding temperature during nitriding becomes a solution to be free from Mg-concentration and related surface oxidation, and to enable plasma nitriding of Al–Cu–Mg alloys. The plasma nitriding experiment was performed at a much lower holding temperature to reduce magnesium diffusion and oxidation. In Table 7, the holding temperature during the presputtering and plasma nitriding processes was fixed at 623 K.

Materials characterization. The over view of the plasma-nitrided AA2014 disc specimen at 623 K for 14.4 ks is shown in Figure 15a. The entire surface is colored black. Its XRD diagram is shown in Figure 15b. Just as analyzed in Figure 9 for the nitrided AA2011 specimen, this diagram is characterized by two peaks for the aluminum matrix and triplet peaks for AlN. This proves that AA2014 alloy is plasma nitrided to form the AlN–AA2014 composite layer via the AlN-cluster nucleation and growth mechanism, in a manner similar to the nitriding behavior observed in the AA2011 specimen. Under the lower holding temperature, the magnesium migration through the Al–Cu–Mg alloy matrix is significantly retarded, sustaining the precipitation reaction to synthesize AlN between aluminum in the matrix and the diffusing nitrogen solutes via the total reaction of Al + N → AlN, with the aid of Al2Cu precipitates.

Figure 15.

RF–DC plasma nitrided AA2014 at 623 K. a) Surface structure of nitrided AA24 specimen, and b) its XRD diagram.

Mechanical characterization. Under the RF–DC plasma nitriding conditions outlined in Table 7, the nitriding duration was reduced to 7.2 ks to describe the surface hardening of AA2014 in the early stage of nitriding. The variation of surface hardness with increasing applied load during micro-Vickers testing is depicted in Figure 16. Although the surface hardness was lower than 800 HV for the RF–DC plasma-nitrided A2011 at 723 K (or 450 °C) for 14.4 ks, the A2014 alloy was also plasma-nitrided to achieve a higher hardness of 400 HV compared to the matrix hardness of 100 HV. This confirms that the nitriding process is effective even in the Al–Cu–Mg ternary alloy system, without requiring Mg-concentration at the surface or causing significant oxidation.

Figure 16.

Variation of the surface hardness with increasing the applied load for RF–DC plasma nitrided AA2014 at 623 K for 7.2ks.

5.3 Application of RF–DC plasma nitriding to aluminum alloy parts

Among a wide variety of aluminum automotive parts, a heatsink has become attractive for the application of duralumin alloys. The heatsink part requires high specific strength and high thermal conductivity to promote the cooling capacity of heated power transistors or LEDs (light emitting diodes). The AlN-capped duralumin alloy heatsink is expected to have high thermal conductance and wear resistance. Most aluminum automotive parts have been made of die-cast or gratitude-cast alloys. In the present experiment, the die-cast AA2011 heat sink was first fabricated to have fin alignment for heat transfer, where the fin thickness was 1 mm, the fin height was 10 mm, and its distancing pitch between adjacent fins was 2 mm, as shown in Figure 17a.

Figure 17.

A die-cast AA2011 aluminum-alloy sink. a) as-die-cast AA2011 heat sink, and b) rf–dc-plasma nitrided AA2011 heat sink at 673 k for 14.4 ks.

Figure 17 compares the plasma-nitrided AA2011 heatsink part at 673 K for 14.4 ks under the conditions in Table 6 to the bare die-cast AA2011 heatsink. The head and side surfaces of the fins are finely nitrided to achieve a surface hardness of 800 HV. This AlN-capping onto AA2011 die-cast heatsink demonstrates the high feasibility of the RF–DC high-density plasma nitriding process in industrial applications.

6. Discussion

In the DC- and RF–DC plasma nitriding of Al-Cu alloy or AA2011 at 673 K, an AlN-rich surface layer is formed as a composite of AlN and Al2Cu within the aluminum alloy matrix. This Al2Cu plays a catalytic role in driving the reaction via Al2Cu + 2 N → 2AlN + Cu and 2Al + Cu → Al2Cu. The nitriding behavior, facilitated by this catalytic reaction, is characterized by the rapid formation of the AlN layer through the overall reaction of Al (in matrix) + N (diffusing nitrogen) → AlN. The isolated nitrogen solutes from NH radicals and nitrogen ions, in addition to activated nitrogen atoms, do not diffuse through AlN but rather through the interfacial boundaries between AlN-clusters and catalytic Al2Cu precipitates. This Al2Cu accommodating both the catalytic reactions and the nitrogen solute diffusion network, drives the growth of a thick AlN-layer at a faster rate than that observed in the normal nitriding process [9–14].

DC-nitriding process of Al-Cu alloys is designed to improve the yield of NH radicals by using a mixture of nitrogen–hydrogen gas with a flow rate ratio of 75% for nitrogen and 25% for hydrogen. This NH radical promptly decomposes into activated nitrogen and hydrogen atoms on the surface of the Al-Cu alloy. The nitrogen atom impinges into the matrix as a reactive solute to Al2Cu. The hydrogen atom recombines to form a hydrogen molecule in the gas phase. RF–DC plasma is ignited under the mixture gas with 80% nitrogen and 20% hydrogen. The yield ratio of the NH-radical population to N2+ population reaches its maximum, driving the nitriding process through chemical nitriding via NH-radical attack on the Al–Cu alloy, together with mechanical nitriding via N2+ bombardment and nitrogen penetration.

Let us discuss the effect of hydrogen flow rate on this yield ratio in the RF–DC plasma nitriding. Figure 18 depicts the variation of maximum surface hardness with an increasing H2/N2 ratio in the mixture gas. When the ratio is less than 20%, or the N2 to H2 ratio is more than 4 to 1, the maximum surface hardness becomes nearly the same, ranging from 950 HV to 1000 HV. A relatively thick AlN-layer is yielded in this low-ratio regime. On the other hand, the maximum surface hardness monotonously decreases with the H2/N2 ratio in Figure 18. This is because of the so-called hydrogen quenching effect [29–31]. This reveals that efficient nitriding could be advanced by RF–DC plasmas under the optimum control of gas supply.

Figure 18.

Variation of measured maximum surface hardness with increasing H2/N2 ratio in the mixture gas.

XRD diagrams in Figures 2 and 9 reveal that AlN formation by nitriding for 14.4 ks becomes equivalent to each other for DC- and RF–DC plasma nitriding.

Magnesium in AA2014 or Al–Cu–Mg alloys poses a risk of deteriorating the nitriding behavior due to its concentration at the surface via migration from the depth and its oxidation at the surface when the holding temperature exceeds 700 K. This deterioration can be suppressed by lowering the holding temperature to 623 K. This reveals that magnesium migration to the surface and oxidation is a thermally activated process; the original AlN formation process can be reproduced by plasma nitriding at lower temperatures. During plasma nitriding, the holding temperature must be lowered as much as possible to suppress the surface diffusion of low melting-point metal elements in the aluminum alloy matrix.

Let us discuss how to effectively lower the holding temperature while preserving the nitriding capacity to form a thick AlN-layer in various aluminum alloys.

As stated in [37, 38], austenitic and martensitic stainless steels were successfully RF–DC plasma nitrided by massive nitrogen supersaturation (MNS). A thick nitrogen-supersaturated layer with a thickness of 60 μm was formed without any nitride precipitates by RF–DC plasma nitriding at 673 K. Due to this MNS, the hardness was enhanced to 1400 HV. As analyzed in [39, 40], the homogeneous MNSed layer was also formed by RF–DC plasma nitriding at 623 K for fine-grained AISI316 and AISI316L stainless steels. This success of low-temperature RF–DC plasma nitriding reveals that massive nitrogen supersaturation provides a way to form a thick solid-solution-hardened layer even at lower temperatures than 600 K without the use of intermetallic compound (IMC) precipitates.

As a typical IMC precipitate in Al–Cu and Al–Cu–Mg alloys, Al2Cu plays a catalytic role in initiating the exothermic reactions to form the AlN layers and sustaining the chained reaction for the growth of the AlN-layer. In the presence of an ignitor to drive the reaction between the diffusing nitrogen solute and the fine IMC-precipitate (IMP), such as Al2Cu precipitate to drive 2 N + Al2Cu → 2AlN + Cu can be initiated. The subsequent reaction via 2Al + Cu → Al2Cu can be sustained even at low holding temperatures. That is, AlxM for x > 1 works as an ignitor to onset the reaction of xN + AlxM → xAlN + M; then, xAl + M → AlxM recovery reaction should take place at lower holding temperatures. In the case of low-temperature plasma nitriding for other series of aluminum alloys such as AA5000, AA6000, or AA7,000, alternative IMCs must be designed to act as catalysts to sustain the exothermic reactions even at temperatures lower than 600 K.

Heat sink made AA2011 alloy was successfully nitrided at 673 K for 14.4 ks, achieving a hardness of more than 800 HV. This homogeneous nitriding in the heat sink fins ensures that duralumin parts and members with small holes and asperities are effectively nitrided using the present approach. This plasma nitriding method offers a new way of surface treatment for various duralumin alloy parts and members. Toward efficient plasma nitriding of both Al–Cu and Al–Cu–Mg alloys in commercial grades, the present RF–DC plasma system can be further improved to enable low-temperature nitriding through high densification of ions and radicals under relatively low electron temperatures.

7. Conclusion

DC- and RF–DC plasma nitriding of Al–Cu alloys at 673 K proves that a thick AlN-rich composite layer is rapidly formed as a hardened surface region. In this nitriding process, Al2Cu plays a catalytic role in initiating and sustaining the exothermic reactions. Fine precipitates in Al-Cu alloys result in the homogeneous formation of the AlN-rich composite layer. In the case of plasma nitriding the Al–Cu–Mg alloys, the holding temperature must be sufficiently low to retard the magnesium solute diffusion to the surface and to minimize the disturbance to the nitriding behavior caused by surface oxidation. Without deterioration due to magnesium diffusion, the AlN-rich layer can be formed for 14.4 ks even at 623 K in a similar manner to DC-plasma nitriding at 673 K.

An A2011 heat sink was prepared by die-casting and RF-DC plasma nitrided at 673 K for 14.4 ks to form a homogeneous, thick AlN-layer. Its high hardness and thermal conductivity make it suitable for heat spreading in power transistors, CPUs, and telecommunication relay stations.

Two approaches toward a lower plasma nitriding process are proposed for further development: massive nitrogen supersaturation into aluminum alloys, and intermetallic precipitate accommodation to AlN-layer formation. In the former, a massively nitrogen-supersaturated layer is formed in aluminum alloys using the fine grain boundary network as a nitrogen diffusion path to sustain the nitriding behavior without the formation of nitrides, including AlN. In the latter, intermetallic compound precipitates (IMP) are designed to work as a catalyst to accommodate the AlN-layer into the aluminum alloy matrix at a lower holding temperature. These more advanced nitriding processes enhance the reusability for upgrading recycled aluminum and aluminum alloys.

Acknowledgments

Authors would like to express their gratitude to Prof. Patama (PhD student at the University of Tokyo; now Associate Professor, Chulalongkorn University, Thailand) and S.–I. Kurozimi (Nano-Coat Lab., LLC.) for their help in experiments.

Conflict of Interest

The authors declare no conflict of interest.

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Written By

Tatsuhiko Aizawa, Hiroki Nakata and Takeshi Nasu

Submitted: 10 September 2025 Reviewed: 25 November 2025 Published: 04 March 2026

© 2026 The Author(s). Licensee IntechOpen. This content is distributed under the terms of the Creative Commons Attribution 4.0 International License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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